ACTAUNIVERSITATIS

UPSALIENSISUPPSALA

2018

Digital Comprehensive Summaries of Uppsala Dissertationsfrom the Faculty of Science and Technology 1740

An X-ray Spectroscopic Study of Perovskites Oxides and Halides for Emerging Devices

DIBYA PHUYAL

ISSN 1651-6214ISBN 978-91-513-0493-9urn:nbn:se:uu:diva-364407

Dissertation presented at Uppsala University to be publicly examined in Häggsalen,Ångströmlaboratoriet, Lägerhyddsvägen 1, Uppsala, Friday, 14 December 2018 at 13:15 forthe degree of Doctor of Philosophy. The examination will be conducted in English. Facultyexaminer: Research Director Manuel Bibes.

AbstractPhuyal, D. 2018. An X-ray Spectroscopic Study of Perovskites Oxides and Halides forEmerging Devices. Digital Comprehensive Summaries of Uppsala Dissertations from theFaculty of Science and Technology 1740. 84 pp. Uppsala: Acta Universitatis Upsaliensis.ISBN 978-91-513-0493-9.

This thesis investigates the electronic structures on several perovskite oxide and halide materialswith a focus on light harvesting applications. The systematic study of the electronic properties ofthe transition metal oxides and post-transition metal halides is a key point if one is to understandtheir properties. The element and site selective nature of several x-ray based spectroscopictechniques are given special emphasis in order to obtain a complete picture of the electronicproperties of the compounds in question. Much of the experimental studies are accompanied byab initio calculations that corroborate with our experimental results.

In the oxide portion of this work, a new class of metallic oxides based on doping of anantiferromagnetic LaFeO3 was synthesized and systematically studied with x-ray absorption,x-ray emission, and photoemission spectroscopies. The compound’s electronic structure iscomplex, having itinerant as well as localized components that give rise to a unique physicalstate where antiferromagnetism, metallicity and charge-disproportionation coexist. Our resonantphotoemission results establish that the Fe states in both magnetically ordered oxides showinsulting properties, while the Mo states provide an itinerant band crossing the Fermi level.An excitation energy-dependent RIXS investigation on LaFe1-xMoxO3 and the double perovskiteSr2FeMoO6 revealed a double peak structure located in proximity to the elastic peak that isidentified to purely d-d excitations, attributed to the strongly correlated nature of these transitionmetal compounds.

The growth of high-quality thin film ferroelectric based on BaTiO3 grown epitaxially bymeans of pulsed laser deposition were investigated. We systematically reduce the band gap ofthe ferroelectric thin film while retaining its polarization at ambient conditions in spite of thealiovalent doping. The electronic structure is studied by several x-ray techniques that show howthe ferroelectricity persists as well as the effective reduction of the band gap through hybridizedstates.

In the post-transition metal halides, the valence and conduction bands were mapped using x-ray absorption, emission, and photoemission spectroscopies. The spectroscopic results identifythe constituent states that form the valence band as well as the band energy positions, which is animperative parameter in optoelectronic devices. In addition, x-ray based spectroscopy was usedto demonstrate the stereochemical activity of lone-pair states (5s2 and 6s2) for several differenthalide compounds and their influence on the chemical, structural, and electronic propertiesof the material. Nanostructured halide perovskites are also explored. The position of iodine pstates and valence band states in reduced dimensional lead-based compounds were examined,as their states are found to be confined in one crystallographic direction in contrast to theirthree-dimensional counterpart. This information highlights the interesting material propertiesand their use in current third generation solar cell research.

Keywords: perovskite oxides, halide perovskites, x-ray spectroscopy, electronic structure

Dibya Phuyal, Department of Physics and Astronomy, Molecular and Condensed MatterPhysics, Box 516, Uppsala University, SE-751 20 Uppsala, Sweden.

© Dibya Phuyal 2018

ISSN 1651-6214ISBN 978-91-513-0493-9urn:nbn:se:uu:diva-364407 (http://urn.kb.se/resolve?urn=urn:nbn:se:uu:diva-364407)

To my family

List of Papers This thesis is based on the following papers, which are referred to in the text by their Roman numerals.

I. Doping induced site-selective Mott insulating phase in LaFeO3

Somnath Jana, Swarup K. Panda, Dibya Phuyal, Banabir Pal, Soham Mukherjee, Anirban Dutta, Puri A. Kumar, Daniel Hed-lund, Joahn Schött, Patrik Thunström, Yaroslav Kvashnina, Håkan Rensmo, M. Venkata Kamalakar, Carlo U. Serge, Peter Svendlindh, Klas Gunnarsson, Silke Biermann, Olle Eriksson, Olof Karis, and D.D. Sarma Submitted, Nature Communications, 2018.

II. Electronic structure of the charge disproportionate and metallici-

ty in LaFe1-xMoxO3 studied by resonant x-ray spectroscopies Dibya Phuyal, Somnath Jana, Soham Mukherjee, Håkan Rensmo, Sergei Butorn, and Olof Karis. In Manuscript

III. Valence Electronic Structure of the double perovskite Sr2FeMoO6 through resonant x-ray spectroscopies Dibya Phuyal, Somnath Jana, Håkan Rensmo, Sergei M. Bu-torin, and Olof Karis In Manuscript

IV. The origin of low bandgap and ferroelectricity of a co-doped

BaTiO3

Dibya Phuyal, Soham Mukherjee, Shyamashis Das, Somnath Jana, Kristina O. Kvashnina, D.D. Sarma, Håkan Rensmo, Ser-gei M. Butorin, and Olof Karis.

Accepted, Europhysics Letters, G40482/B19376, 2018. V. Ferroelectric properties of doped BaTiO3 thin film by pulsed la-

ser deposition Dibya Phuyal, Soham Mukherjee, Somnath Jana, Shyamashis Das, D.D. Sarma, Håkan Rensmo, Sergei M. Butorin, and Olof Karis. In Manuscript

VI. Impact of synthetic routes on the structural and physical proper-

ties of butyl-1,4-diammonium lead iodide semiconductors

Majid Safdari , Dibya Phuyal , Bertrand Philippe, Per H. Svens-son, Sergei M. Butorin, Kristina O. Kvashnina, Håkan Rensmo, Lars Kloor, and James M. Gardner Journal of Materials Chemistry A, 5, 11730, 2017.

VII. Electronic Structure of Two-Dimensional Lead(II) Iodide Perov-

skite: An Experimental and theoretical Study Dibya Phuyal, Majid Safdari, Meysam Pazoki, Peng Liu, Ber-trand Philippe, Kristina O. Kvashnina, Olof Karis, Sergei M. Butorin, Håkan Rensmo, Tomas Edvinsson, Lars Kloo, and James Gardner Chemistry of Materials, 30, 4959-4967, 2018.

VIII. The electronic structure and band interface of cesium bismuth

iodide on a titania heterostructure using hard X-ray spectroscopy Dibya Phuyal, Sagar M. Jain, Bertrand Philippe, Malin B. Jo-hansson, Meysam Pazoki, Jolla Kullgren, Kristina O. Kvashni-na, Matthias Klintenberg, Erik M. J. Johansson, Sergei M. Bu-torin, Olof Karis, and Håkan Rensmo. Journal of Materials Chemistry A, 6, 9498, 2018.

IX. An effective approach of vapour assisted morphological tailor-ing for reducing metal defect sites in lead-free, (CH3NH3)3Bi2I9 bismuth-based perovskite solar cells for improved performance and long-term stability Sagar M. Jain, Dibyua Phuyal, Matthew L. Davies, Meng Li, Bertrand Philippe, Catherine De Castro, Zhen Qiu, Jinhyun Kim, Trystan Watson, Wing Chung Tsoi, Olof Karis, Håkan Rensmo, Gerrit Boschloo, Tomas Edvinsson, and James R. Durant. Nano Energy, 49, 614-624, 2018.

X. Inorganic CsPbI3 Perovskite Coating on PbS Quantum Dot for Highly Efficient and Stable Infrared Light Converting Solar Cells Xiaoliang Zhang, Jindan Zhang, Dibya Phuyal, Juan Du, Lei Tian, Viktor A. Öberg, Malin B. Johansson, Ute B. Cappel, Olof Karis, Jianhua Liu, Håkan Rensmo, Gerrit Boschloo, and Erik M.J. Johansson. Advanced Energy Materials, 8, 1702049, 2018.

Extended Bibliography The following is a list of publications to which I have contributed, but not included in this thesis.

1. Two-Dimensional flexible high diffusive spin circuits Ismael Serrano, Jaganandha Panda, Fernand Denoel, Orjan Val-lin, Dibya Phuyal, Olof Karis, Venkata M. Kamalakar. Submitted, Nano Letters, Aug 2018

2. Investigation of the spectral properties and magnetism of BiFeO3 by dynamical mean-field theory Souvik Paul, Diana Iusan, Patrik Thunström, Yaroslav O. Kvashnin, Johan Hellsvik, Manuel Pereiro, Anna Delin, Ronny Knut, Dibya Phuyal, Andreas Lindblad, Olof Karis, Biplap Sanyal, Olle Eriksson Physical Review B 97, 125120 2018.

3. Cesium bismuth iodine, CsxBiyIz, solar cell compounds from systematic molar ratio variation Malin B. Johansson, Bertrand Philippe, Amitava Banerjee, Dib-ya Phuyal, Sudip Chakraborty, Mathis Cameau, Huimin Zhu, Rajeev Ahuja, Gerrit Boschloo, Håkan Rensmo, and Erik M.J. Johansson Submitted, ACS Material and Interfaces, 2018

4. Exchange dependent ultrafast magnetization dynamics in Fe1-

xNix alloys Somnath Jana, Ronny Knut, Erna K. Delczeg-Czirjak, Rameez S. Malik, Robert Stefanuik, Raghuveer Chimata, Dibya Phual, Venkata Mutta, Serkan Akansel, Daniel Primetzhofer, Martina Ahlberg, Johan Söderstrom, Joahn Åkerman, Peter Svedlindh, Olle Eriksson, and Olof Karis In Manuscript

Comments on my own participation:

The studies presented in this thesis are the result of collaborations between experimentalists and theoreticians. I have been the responsible person for preparing and preforming x-ray photoemission as well x-ray absorption and emission spectroscopy measurements, treating the experimental data, and writing these relevant sections for all manuscripts. I have been the primary responsible person in fabricating and characterization of thin films as well as spectroscopic results for papers IV and V.

Reprints were made with permission from respective publishers.

Contents

1 Introduction ................................................................................................ 111.1 General Overview: Emerging Materials ......................................... 111.2 Properties of Perovskite Oxides ......................................................... 13

1.2.1 ABO3 Perovskite Properties and Electronic Structure ................ 131.2.2 Ferroelectric Oxides .................................................................... 16

1.3 Inorganic-Organic Halide Perovskites ................................................ 181.3.1 ABX3 Crystal and Electronic Structure ...................................... 191.3.3 Solar Cell Devices and Interfaces ............................................... 22

1.4 Aim of Thesis ..................................................................................... 24

2 X-Ray Spectroscopic Techniques and Instrumentation ............................. 252.1 Basics of X-Ray Spectroscopy ........................................................... 25

2.1.1 Synchrotron Radiation: Tunable Soft & Hard X-rays ................ 272.1.2 Storage Rings and Insertion Devices .......................................... 272.1.3 Beamlines and Endstation ........................................................... 28

2.2 Principles of Photoemission Spectroscopy (PES) .............................. 292.2.1 General Formalism: The Photoemission Process ....................... 292.2.2 Electron Escape Depth and Photoionization ............................... 312.2.3 Resonant Photoemission Spectroscopy ....................................... 31

2.3 Principles of X-Ray Absorption and Emission Spectroscopy ............ 332.3.1 Introduction to XAS and HERFD-XAS ..................................... 332.2.2 X-Ray Emission Spectroscopy ................................................... 352.2.3 Resonant Inelastic X-Ray Scattering .......................................... 37

2.4 Thin Film Fabrication of Ferroelectrics .............................................. 392.4.1 Pulsed Laser Deposition ............................................................. 392.4.2 X-ray Diffraction and Reflectivity .............................................. 402.4.3 Piezoforce Microscopy ............................................................... 40

3 Bulk Perovskite Oxides .............................................................................. 423.1 Insulator-to-Metallic Antiferromagnetic Oxide: LaFe1-xMoxO3 ......... 423.2 Valence Electronic Structure of LaFe1-xMoxO3 .................................. 453.3 The Double Perovskite Sr2FeMoO6 .................................................... 49

4 Bulk and Thin Films of a Doped BaTiO3 ................................................... 534.1 Low Bandgap and Ferroelectricity of co-doped BaTiO3 .................... 534.2 Thin Film Properties of co-doped BaTiO3 ......................................... 55

5 Halide Perovskites ...................................................................................... 595.1 Low-Dimensional Halides .................................................................. 595.2 Lead-Alternative Halide Perovskites .................................................. 635.4 Nanocrystal Halides ............................................................................ 68

6 Outlook ....................................................................................................... 70

7 Acknowledgements .................................................................................... 71

8 Populärvetenskaplig Sammanfattning ........................................................ 73

9 References .................................................................................................. 76

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1 Introduction

1.1 General Overview: Emerging Materials

This thesis targets the development and understanding of emerging materials for fundamental meaning and their connection to early-stage technologies. My objective is to develop an understanding of the electronic and/or magnet-ic behavior of complex oxides and halide perovskites. This understanding is an essential step to develop new functionalities that can benefit electronic and energy related applications. It is broadly recognized that the era of microelectronic devices faces many hurdles if it is to meet its expectations outlined by ‘Moore’s Law’ stated in the 1970s1. The already exceptional performance of modern electronic de-vices has transformed much of how society uses and interacts with technolo-gy. However, the multitude of scaling limitations with CMOS-based tech-nology and prolific use of communication technology infrastructure necessi-tates a new era for new electronics. It is predicted that the information and communication sectors and memory electronics will consume as much 40% of global electricity supply by 20302. Therefore, the need for alternative technologies must be envisaged to meet this demand3,4. Finding new material systems that can be implemented at a scalable cost, meet the material availa-bility, and integrate into existing manufacturing process is the incentive for much of the academic and industrial research into complex metal oxides today. Complex oxides have been one of the main pillars of solid-state physics over the last century and continue to offer novel physics. Recent work in oxide electronics suggests it may be plausible to implement such systems at the electronic device level, with potential to reshape computational density and energy-efficiency5. This wide range of materials is expanding beyond the scope of integrated electronics into applications previously not emphasized such as photovoltaics6,7, photocatalysis8, water splitting9, and even into health related applications10. Advances in thin film deposition, fabrication of heterostructures, and characterization has ushered a new era of complex oxides to that can be harnessed for multi-functional devices.

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The other material class investigated in this thesis is perovskite halides for energy-related applications. Halide perovskites have emerged as a serious contender in this field11,12. Although silicon and thin film technologies costs have significantly reduced in the last two decades, their fabrication process and materials input to manufacture can still benefit from improvement13. Finding alternative materials, with improved performance and simplified engineering, has incentivized the search for new materials that can lower the cost even further. This material class, whose nomenclature is partly adopted from the extensively studied perovskite oxide family, shows exceptional material properties and device performance that is also associated with low-cost solution based fabrication process analogous to the successfully com-mercialized organic light-emitted diode televisions14. The rapid advancement of this material is usually linked to optimization on the elemental composi-tion, high crystal quality, and minimizing defects by achieving sharper inter-faces at film junction, and device architecture. However, several issues need to be resolved to validate this material for commercial relevance. For instance, the stability issues related in the prevail-ing halide perovskites has remained a concern, as well as the environmental and toxicology hazards of lead15–18. This has led to innovations in finding more suitable materials through compositional changes without comprising the outstanding properties of lead-based halides. One route has been exploit-ing the crystal structure to form lower dimensional variants to the halide perovskites, which has surpassed stability performance of the highest effi-ciency three-dimensional variety19. The other route is the replacement of Pb ions with alternative metals that can retain much of the optical and electrical properties of the best-preforming lead-based materials20–22. Since it is established that the electronic structure of perovskite halides are directly linked to some key properties23,24, this thesis investigates the elec-tronic structures of two emerging routes in halide perovskites: reduced-dimensional systems and lead-alternative materials. Fundamentally, the development of these emerging materials typically starts from the elemental and atomistic description of the constituent material. The atomic-scale engineering of exotic new materials is mostly of fundamental interest, however the information obtained could serve concrete interest for materials development. A systematic study on the electronic structure of materials such as transition metal oxides and perovskite halides is a key point in understanding their macroscopic properties. This thesis focuses on these challenges in materials research by using x-ray spectroscopic methods. The information from x-ray spectroscopic techniques can offer clues into their electronic structure and underlying chemical and physical mechanisms. In this way, the knowledge obtained from x-ray spectroscopy is indispensi-

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ble in correlating the electronic structure to the macroscopic properties of emerging materials.

In this chapter, a brief state of the art review about the emerging materials in oxide research and inorganic-organic halides are presented.

1.2 Properties of Perovskite Oxides

1.2.1 ABO3 Perovskite Properties and Electronic Structure

The bulk transition-metal oxide (TMOs) perovskites have displayed a wide range of phenomena that have generated special interest to experimentalist and theorists alike. Although this class of materials are well-studied25, their intriguing properties such metal-insulator transitions, colossal magnetore-sistance, ferroelectricity, and superconductivity continue to surprise the re-search community26–29. These properties emerge from interactions of elec-trons in a many-particle system that is principally governed by several com-peting interactions: charge, spin, lattice, and orbital as shown in Fig. 1.1(a). The most significant of these interactions is the correlation of electrons in perovskite solids, where the interplay between electrons originate from hy-bridization and localized effects often rooted in electron-electron repulsions27. This central parameter to perovskite oxides, the interaction of the 3d electrons, can induce ground states with unusual macroscopic charac-teristics.

Figure 1.1 (a) The competing interactions of charge, spin, lattice, and orbital de-grees of freedom in transition metal oxides (Figure adapted from Ref. [29]) (b)Zaanen-Sawatzky-Allen framework successfully applied to most TMOs based ontheir U, W, and Δ parameters.

Figure 1 1 (a) The competing interactions of charge spin lattice and orbital de

�<U'

U'

U'

dr

U'<� U'<�

W

14

The metallic or insulating behavior of oxides has launched systematic stud-ies on the many interactions and correlation in solids. In the standard para-digm of perovskite oxides, the global electronic structure is understood through the competition and/or cooperation by the TM 3d-3d on-site Cou-lomb interaction energy known as U, the charge-transfer energy from the occupied O 2p orbitals to the empty TM 3d orbitals denoted by Δ, and the bandwidth W. A successful model that identified correlated systems into separate class of materials of many TMOs is the Zaanen-Sawatzky-Allen framework30, which recognized the critical parameters that predict whether oxides are metallic or insulating31 as shown on the right panel of Fig 1.1(b). Strong interplay between the bandwidth (W), electron-electron repulsion (U), and the charge transfer (Δ) between metal and oxygen bands reveal separate classes of materials reflected in these values:

• Mott-Hubbard Insulators: The band gap is proportional to U in the case

U Δ, where the Coulomb process dominates and reveals an energy gap between inter-site transition metal ion.

• Charge Transfer Insulator: This type is among the most common to perovskite oxides, where the band gap is proportional to Δ, in which the lowest excitation is transferring an electron from the O ligand states to the transition metal ion.

• Intermediate Compounds: If Δ ~ U, where a competition between both Coulomb and charge transfer excitations exists.

• Metals: This class can be divided into two sets: a semi-metal (Δ < W), and the other a soft metal (U < W).

Based on this model, it is possible to describe the variation of the band gap of insulators on U and Δ. These properties also are dependent on the compo-sition through doping, as well as long-range changes to the crystal structure through stress, temperature, and external field. This alters the electronic phase diagram and the ground state properties of the comound. A demonstra-tion of this effect is the colossal magnetoresistance observed in La1-

xSrxMnO3 by doping the antiferromagnetic (AFM) charge-transfer insulator LaMnO3

32. This electronic reconstruction is associated with a half metallic ferromagnetism that is explained in terms of charge carriers and local spins from Mn 3d states33. This example is one of many in which a collective phase transition through doping enables oxide materials for interesting fun-damentals and for potential use in emerging devices.

As discussed earlier, perovskite oxides have bands that are strongly associat-ed with d orbital character. A transition metal (TM) in a crystal structure with octahedral (or several other symmetries) symmetry leads to lifting the degeneracy of d orbitals due to the electrostatic interaction with neighboring O 2p orbitals. The interaction of a TM ion in a ligand electrostatic field from

15

a quantum mechanical approach is called Crystal Field Theory (CFT). The orbital quantum number for electrons in a d orbital is L = 2, thus the degen-eracy is 2L+1 = 5. These orbitals are: dxy, dyz, dzx, dx2- y2, d3z2- r2, and are illus-trated in Fig. 1.2(a). These orbitals are also distributed preferentially in space and have their distribution maxima in a region in the coordinate axes, de-pending on the orientation of the lobes. In the octahedral environment, they can be classified in two groups: dxy, dyz, and dzx, which point in between the coordinate axes are known as t2g, and dx2- y2, d3z2- r2 have a maxima along the coordinate axes, known as eg orbitals.

When the central TM ion is placed in a non-spherically symmetric electro-static field, the different d orbitals respond non-equivalently. The energy of the individual d orbital adjusts to this crystal field and shifts according to its orientation in the unit cell. An example of this effect on a TM ion in a octa-hedron coordination, where the eg orbitals point toward the oxygen ligands increasing their energy, while the t2g orbitals point in between them, result-ing in a lower energy for the latter as shown in Fig. 1.2(b). The residual de-generacy can be lifted even further when considering the coordination (such as tetrahedral, orthorhombic) of the TM ion and the different distortions of the octahedra. The electron filling of these orbitals are complicated further by other interactions such as spin-orbit coupling (LS), Hund’s coupling (JH), Jahn-Teller distortions, or other coupled phenomena and beyond the scope of this thesis.

The spectroscopic signatures of electronic phase transitions and the crystal field splitting are explored in this work through several x-ray spectroscopic methods. Photoemission spectroscopy, in particular at the resonant condi-tion, is perfectly suited to qualitatively asses the changes to the electronic

Figure 1.2 (a) The five-fold degenerate d orbitals in a transition metal ion (Figure adapted from Ref. [27]). (b) The crystal field splitting lifts the degeneracy of the d orbitals for different symmetries.

zx yz xy

3z2-r2 x2-y2

eg Orbitals

t2g Orbitals

t2g

Free Atom Cubic(Octahedron)

Tetragonal Orthorhombic

eg

4

10

2

2

4

2

2

2

22

2 dxy

dzx

dyz

d3z2-r2

d3x2-y2

6

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ground state of a Mo doped LaFeO3 where we discuss the transition from an insulating to a metallic state as a function of doping. The double perovskite Sr2FeMoO6, where Fe and Mo occupy alternatively at the B site of the unit cell, is also rationalized through x-ray spectroscopy on its unique coexist-ence of half-metallic and magnetism. The material properties and the de-tailed electronic structure results are presented in Chapter 3.

1.2.2 Ferroelectric Oxides

Ferroelectric materials possess a spontaneous electrical polarization that can be switched with an external electrical field. This stable ground state is due to a minimized free energy configuration with a double potential well that is the result of a phase transition from an ordinary dielectric to ferroelectric at the Curie temperature (TC) as shown in Fig. 1.3(a). The most distinguishing feature of ferroelectrics is the possibility to switch the spontaneous polariza-tion by an applied field. The spontaneous polarization of ferroelectric mate-rials implies a hysteresis effect, which can be used as a memory function34, shown in Fig. 1.3(b). It is noteworthy that the FE polarization has a remnant polarization (PR) after the sweeping of electrical voltage, similar to the mag-netic analogue in ferromagnetic materials.

BaTiO3 (BTO) is a prototypical ferroelectric oxide and is a special interest for its room temperature polarization and used in many applications such as data storage, photovoltaics, transducers and actuators among others35. The phase transition is from a high-temperature cubic dielectric to a te-tragonal ferroelectric at T ~ 400 K, and orthorhombic to rhombohedral at 190K. The origin in the phase transition from a cubic structure dielectric to a non-centrosymmetric tetragonal phase at TC is accompanied by a off-

T < TC

T = TC

T > TC

P

F

Figure 1.3 (a) The free energy (F) is plotted as function of polarization for T >Tc, T = Tc, and T < Tc (b) Schematic of a polarization hysteresis loop for a typi-cal ferroelectric in a coercive field (EC), the spontaneous polarization, PS, andthe remnant polarization, PR. (c) The off-centering of the Ti ion in BTO, wherethe polarization P points in the direction of the ion displacement.

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centering of the smaller Ti cation, where the spontaneous polarization de-rives largely from the electrical dipole moment created by this shift36. The resulting polarization is directed along the c axis, in the same direction as the Ti ion displacement (see Fig. 1.3(c)). This phase shift is accompanied by changes to the local structure and overall covalency, where the significant hybridization between Ti 3d states and O 2p states is essential in stabilizing the ferroelectric distortion36. The lowest energy transition in the charge transfer-type insulator is from the strongly mixed Ti d - O 2p states to unoc-cupied Ti 3d bands with a gap (Eg) of 3.2 eV. The Ti d0 configuration is rea-soned to be a necessary condition for B cation displacement, as other BaMO3 (M = Mn, Fe, Co, etc.) oxides with formally occupied dn states at the metal site retain their inversion symmetry37. The formal valency at the d site in is also concomitant with magnetic order, as d orbital occupancy is requirement for magnetic moments. However, the d0 and dn formal valency tend to com-pete for ferroelectric and ferromagnetic order, with only a few room temper-ature materials with a coupled ordering37. Another extraordinary property of ferroelectric oxides is the ability to gener-ate a bulk photovoltaic effect (BPVE)38. Under illumination, they exhibit a photocurrent that is parallel to the polarization direction. Although ferroelec-trics usually have a large optical gap, they have recently shown remarkable results that have attracted the attention of the photovoltaic community. Re-cent experiments showed that a Cr doped BiFeO3 (a room-temperature ferro-electric with weak magnetism) in a planar device architecture under illumi-nation showed an efficiency of 8.1 %7. The BPVE in ferroelectrics is also able to generate a photovoltage that is larger than the bandgap39. BTO was shown to produce a photovoltage of 8.1 V and efficiency of 4.1%40. There is a serious push to improve upon this work with low bandgap ferroelectrics, improved thin film deposition, and domain structure manipulation. This groundbreaking work has pushed the field of photoferroics into new focus and attention of broader research disciplines. In this work, we investigate the effects of compositional changes to the over-all electronic structure on a co-doped BTO discussed in Chapter 4. The in-tentions are to reduce the overall change transfer gap by mixed valence dop-ing at the B site of BaTiO3 with a d0 and dn metal cations respectively. The substitution at the B site and its effects on its electronic structure are ex-plored with complementary x-ray techniques to reveal the nature of bandgap tuning and effects on polarization.

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1.3 Inorganic-Organic Halide Perovskites

Inorganic-organic hybrid perovskites halides (HPs) have emerged as a lead-ing contender in the field of low-cost high efficiency solar cell materials. The key attributes to their emergence in solar cell technology is arguably due to its strong optical absorption, long-range electron-hole diffusion lengths, and reasonable mobility of charge carriers, culminating to a high photovolta-ic efficiency - all achieved within the past decade12,41,42. The primary com-pound in this large class of materials is MAPbI3 (MA = CH3NH3), which has a low bandgap of 1.55 eV and a notable efficiency of 12.1% during its early stages of development in 201243. The efficiencies since then have risen dra-matically to above 22% in a chemically-substituted compound44 that now qualify this material for its place in the market with established technologies like silicon and thin film technologies (like CdTe and CIGS). A diverse range of synthesis methods has been used to fabricate devices, most com-monly using low-temperature methods. These materials have also attracted attention for its cluster of fundamental scientific discoveries that are ad-dressed in several reviews12, 45–47.

The efficacy of these materials in solar cell devices has given ground to fur-ther optimizing this material for commercialization. This has also led to find-

Figure 1.4 The exceptional structural flexibility of hybrid perovskites enables the formation of various types of structures that strongly depend on both compo-sition and synthesis procedures (a) prototypical 3D ABX3 cubic structure (b) 2Dlayered structures separate the inorganic MX6 octahedron, (c) 1D nanowires in which metal-halide octahedrons are connected in a chain, and (d) nanocrystalswhere metal-halide octahedrons are completely isolated (Figure adapted from Ref. [47].)

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ing alternative variations to inorganic-organic perovskites to achieve compa-rable materials for high efficiency photovoltaics, such as exploiting the tre-mendous flexibility to obtain various types of crystal structure as shown in Fig. 1.4. A promising trend is in layered HPs, colloidal nanocrystals, and lead-alternative perovskites – which are all material systems investigated in this thesis. This subset of HPs has unique material characteristics and good device performance of their own, and could individually find their place in future optoelectronic devices.

The subsequent section reviews the chemical and electronic structure of HPs and its different variants. The bulk electronic structure is reviewed and a significant portion of the experimental work in this thesis. The common de-vice architectures in solar cell devices and the role of interfaces are also briefly addressed.

1.3.1 ABX3 Crystal and Electronic Structure

Hybrid organic-inorganic perovskites take their nomenclature from its crys-tal structure that is similar to the CaTiO3 oxide perovskite (ABO3). The most prominent and well studied material in this class is MAPbI3, shown on the far right in Fig. 1.5 A number of exciting physical properties, like colossal magnetoresistance, ferroelectricity and superconductivity, have all been dis-covered in this prolific family of compounds during the past century48. For the ‘hybrid’ perovskites, ones with both inorganic and organic cations, are usually either methylammonium (MA+ = CH3NH3

+) or formamidinium (FA+ = CH(NH2)2

+), large monovalent cations that fill the octahedral interstitials. The all-inorganic compounds usually contain cesium or recently even rubid-ium (Rb) at the A site. The B site is usually a post-transition metal (Pb, Bi,

Figure 1.5 As the tolerance factor deviates from the ideal t = 1, the size re-striction on different cations and anions are lifted to form two-dimensional de-rivatives.

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Sb, or Sn) and located at the center of the body of an octahedral cluster. The octahedral, [PbI6]4- unit can form three-dimensional (3-D), two-dimensional (2-D), one-dimensional (1-D) or zero-dimensional (0-D) networks (see Fig. 1.4). The anion component is usually a halide element (I, Br, or Cl), making the combinatorial possibilities of this compound numerous. This versatility in this class of hybrid materials allows a vast space to explore new materials with new and useful properties.

The ideal cubic structure for perovskite halides (AMX3) has 6-fold coordina-tion within the octahedron and the A cation in a 12-fold cuboctahedral coor-dination. The crystallographic stability and probable structure is empirically deduced by considering the well-established Goldschmidt’s tolerance factor49, a ratio of the interatomic distance in an idealized solid-sphere,

. The slight differences in the A-X and M-X ionic radii are stringent and a large mismatch can introduce buckling and distortions to the equilibrium structure50. This generally produces a lower-symmetry dis-torted version of the ideal cubic structure, which in turn can manifest a varie-ty of structures with a different set coordination numbers to the A and B cations. Aside from the constraints on the ionic radii, the charge balance and valence assignment can further be extended to attain new polymorphs50.

Given that there are chemical and physical constraints to achieve a 3D struc-ture, the substitution of various cation and anion sizes allows for remarkable structural tunability. This can result in deviations in the octahedral size and tilt, as well reduced dimensionality mediated by inorganic layers separated by a cation. In two-dimensional perovskites, where size restrictions on the A site cation is relaxed, can result in structures with a different layering of the inorganic layer in between the organic chain. The charges in two-dimensional (2D) materials are restricted in their movement to the two-dimensional plane; however, their confined structure allows one to tune the optical and electronic properties by varying their thickness (see Fig. 1.5). This is usually given an index number to the stoichiometry of the material, R2(CH3NH3)n-1MnX3n+1, where R is an organic group, M the divalent transi-tion metal, X is a halide, and n refers to the number of metal halide layers between the long chain organic cation layers. This structure converges to a 3D structure with , where the ending compound narrows the bandgap and/or increases the mobility of charges within the layers51. Although their efficiencies have not matched the 3D equivalents, their improvement is on-going and efficiencies of 15.3% were reported using n = 3 inorganic sheets separated by an organic cation1919. The major advantage of 2D HPs over the 3D counterparts is their stability to moisture, which is a serious impediment for successful commercialization of 3D HPs.

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The basic electronic structure of the simplest and well-studied perovskite MAPbI3 is partly understood23,52. Early studies using density functional theo-ry (DFT) calculations showed the inorganic PbX6 octahedra to have an elec-tronic configuration based on Pb: 5d106s26p0 and I: 5p6. The Pb 6s/6p-I 5p interactions are responsible for the chemical bonding and form the valence and conduction bands as shown in Fig. 1.6. Arguably the most important property of a light absorbing material is the bandgap (Eg), which is directly related to the interaction of frontier electronic states of the constituent com-pounds. For HPs, the valence band maximum (VBM) is formed of an anti-bonding (σ*) Pb 6s – I 5p interactions, while the conduction band minimum is by empty Pb 6p and non-bonding Pb 6p - I 5p interactions, as shown in Fig. 1.6(a). The molecular states stemming from the organic or inorganic A cation lie deeper in the VB and do not hybridize with the inorganic units (PbI6 for example) and lack any significant electronic interaction. Although they do not contribute directly to band edge states, they do influence the crystal structure framework, including the symmetry and bond lengths of the PbI6 octahedra and stability of the HPs thin films. This configuration of the electronic structure has led HPs to obtain strong optoelectronic properties including characteristic known as defect toler-ance21. One of the proposed mechanism by which defect tolerance occurs is the presence of antibonding interactions at the VBM, which tends to confine defects to shallow states at the band edges or outside the gap rather than deep in-gap states that act as recombination centers21, as shown in Fig. 1.6(b). Several authors have argued that the key to properties such as defect-tolerant solar absorbers is due to the bonding structure of the lone pair ns2 hybridization at the B site of HPs crystal structure53. In the perovskite struc-

Figure 1.6 (a) Electronic structure and bonding structure common to most per-ovskite halide compounds. (b) The antibonding states at the top of the valenceband maximum give rise to shallower defects.

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ture, a stereochemically active lone pair (or s-orbital hybridization) generally gives rise to structural distortions such as bond lengths and bond angles among other properties54. Materials currently of interest with this configura-tion are Pb2+, Sn2+, Ge2+, Sb3+, and Bi3+, a few of which are discussed in the results section of this thesis.

Although lead-based perovskites have demonstrated the best performance so far, its toxicity and stability present a significant challenge to commercializa-tion. This obstacle has spurred theoretical and experimental research into lead alternatives as a path to overcome the issues pertaining to lead that can deliver comparable material and devices properties that improve upon the toxicity and stability. A few different compositions are explored in their electronic structure in relation to the MAPbI3 analogue, namely with substi-tutions with the trivalent Bi3+ and Sb3+ halides that also have a similar lone pair configuration (ns2).

The final class of HPs investigated in thesis takes the form of nanometer-sized nanocrystals (or quantum dots). Recently, they have attracted a diverse set of material science research with promising results in photovoltaic appli-cations, with photovoltaic efficiencies reaching 10.1 %55. In general, nano-crystals less than 20 nm are quantum-confined in all three directions, and the improved photophysical properties are accredited to nanoscale effects from quantum confinement56. The perovskite quantum dots have garnered atten-tion due to its excellent photophysical properties, bandgap tuning, and high quantum yield of luminescence57. Due to the favorable lattice matching be-tween metal halide perovskite and the cubic rock-salt PbS, perovskites have emerged as a preferred material for surface passivation forming a core-shell type structure58. CsPbI3 nanocrystals coated directly on PbS quantum dots and its surface electronic structure using photoelectron spectroscopy is pre-sented in Chapter 5.

1.3.3 Solar Cell Devices and Interfaces

The most common solar cell architecture for HPs materials are based on the classical dye-sensitized solar cell design with carefully selected layers for efficient charge generation and collection as shown in Fig. 1.7. The princi-ples of operation of this device involve many competing factors and readers are suggested to several reviews for an in-depth analysis11,59,60. In brief, the operation involves the (1) the generation of electron-hole pairs (excitons) in the HP absorbing layer, (2) injection and separation of electron-hole pairs into selective n-type electron acceptors and a p-type hole-transport layers (HTM), (3) and subsequent diffuse of carriers through transport layers to electrodes where they simultaneously generate a voltage and current to the

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external circuit. This simple heterojunction can be adapted to a broad range of materials and fabrication techniques. The promise in this device architec-ture is the prospect for low manufacturing costs associated with an all-solution based fabrication.

A key function in device operation is the efficient separation and transfer of electron-hole charge carriers at the interfaces in adjacent layers. A crucial requirement is the matching of electronic energy levels between the perov-skite absorbing layer and adjacent n-type and p-type layers with minimal energy loss. This should facilitate the efficient and faster charge separation of charges to avoid recombination at the interface, which dramatically hin-ders the efficiency of the cell. The interface energetics is also crucial to de-termining the limitation in the open-circuit voltage60.

The valence and conduction band maximum (VBM and CBM respectively) positions of the absorbing and electron and/or hole conducting layers are shown in Fig. 1.7(b). The n-type electron accepting layer CBM has be to higher in energy, i.e. the band offset, to facilitate efficient separation and injection of electrons and lower in energy for holes to the HTL. This energy scheme can be appropriately identified with x-ray photoelectron experi-ments, and the obtained information can help guide the design of high per-formance structure. The thesis presents quantitative information on the inter-facial energetics for several different perovskite halide systems discussed in the previous section.

Figure 1.7 (a) The most commonly used planar heterojunction architecture where the photoactive HP layer is adjacent to selected electron and hole transport layers (light blue and tan respectively). (b) The favorable kinetics necessary to driveefficient electron-hole carriers to reach the electron transport layer and hole transport layer. (Figure adapted from Ref. [11]).

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1.4 Aim of Thesis

The aim of this thesis is to develop an understanding of the electronic struc-ture of these emerging materials through several x-ray spectroscopic tech-niques. The electronic structures can elucidate much of the observed proper-ties seen in halides and oxide perovskites and guide in materials design. The aim of this chapter has been to introduce the most essential concepts into the materials and the motivation in their research.

Chapter 2 presents the experimental techniques used in this thesis and their abridged formalism. Chapter 3 presents the development of a new material with a rare combination of magnetic order and metallicity, and is demon-strated with extensive electronic structure evidence. Chapter 4 focuses on the development a new low bandgap ferroelectric and how it attains its proper-ties, which is established through various x-ray spectroscopic techniques. The fabrication of thin film and its properties are also presented. Lastly, Chapter 5 discusses the results on halide perovskites supported with theoret-ical calculations.

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2 X-Ray Spectroscopic Techniques and Instrumentation

In this chapter, I will present the different experimental methods used to conduct x-ray spectroscopy measurements. The basic formalism of x-ray spectroscopy is briefly reviewed. This section also summarizes the thin film fabrication by pulsed laser deposition and thin film characterization tech-niques.

2.1 Basics of X-Ray Spectroscopy

The aim of x-ray spectroscopy is to obtain information about the electronic structure of a system. The complexity of x-ray interaction with matter offers many spectroscopic routes to bear information to the materials electronic structure, crystal structure and spatial ordering, magnetic ordering, and much more. Since it is possible to obtain chemical and physical properties of the sample through the interaction of light and the electrons in the host sample, x-ray spectroscopy is one of the most powerful tools in modern physics and chemistry.

In solids, electrons can be separated into two categories: core electrons (in-ner shell) and valence electrons (outer shell). The valence electrons include the outermost and least bound electrons that are responsible for chemical bonding. Its spatial and energetic separation into occupied and unoccupied states primarily account for the observed physical properties of the material; that is electric, magnetic, thermal, and other properties determined by va-lence electrons. The core electrons are deeply bound and are almost un-changed to that of free atoms and are well separated in energy and unique to each element. Although generally inert due to their localization near the nu-cleus of an atom, core electrons provide a method to locally study the va-lence electronic structure and geometry around a specific atomic site61. When an incident photon with sufficient energy interacts with a multi-electron atomic system, an electron can be removed (excited or ionized) from deep electron shells by the photoelectric effect which forms the basis of

26

x-ray absorption spectroscopy (XAS), x-ray photoemission spectroscopy (XPS), and x-ray emission spectroscopy (XES). An electron can absorb a particular energy and is then promoted to a partially filled or empty lev-els/bands just below a particular ionization threshold, schematically illustrat-ed in Fig. 2.1(a). The perturbation to the ground state involves electrons making transitions into new bound states or unbound continuum states. The character of this excited state depends critically on the energy of the incident radiation. The excited state is unstable, and the system will decay after a certain lifetime via emission of photons (radiative decay) or electrons (non-radiative (Auger) decay) shown in Fig. 2.1(b). Measuring the energy and momentum of the emitted particles allows one to deduce various properties of the system in its ground or excited states.

X-ray spectroscopic techniques can be roughly divided into two types of scattering events, one that is a ‘first-order’ scattering event than includes only one photon with incident energy ωi, such is the case for XPS and XAS. The core electron is promoted to an empty valence state by an incident x-ray photon of energy. For XPS, we have a photon absorbed and an electron ejected from the sample. These processes are followed by subsequent relaxa-tion process shown schematically in Fig. 2.1(b). For ‘second-order’ scatter-ing event, a photon promotes an electron to an excited state and decays radi-atively with the emission of an x-ray photon, as is the case for XES and RIXS. There are also three regimes in which x-ray spectroscopy can be di-vided into based on photon energy: soft x-ray region (~ 20 – 2000 eV), ten-

Figure 2.1 (a) Excitation of a core-electron into either a bound state or a continuum state. (b) The relaxation channels after tran-sition to an excited state.

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der x-ray (~ 2000 – 6000 eV), and hard x-ray (~ 6000 – 30000 eV). All three ranges have their own specific qualities as well as offering complimentary information on the overall electronic structure and will be given a brief overview in subsequent sections.

2.1.1 Synchrotron Radiation: Tunable Soft & Hard X-rays

In order to perform the many different x-ray spectroscopy experiments, one needs high intensity and high flux tunable x-ray sources. The state-of-the-art instrumentation for these experiments requires synchrotron radiation (SR). When charged particles moving at relativistic speeds along curved trajecto-ries by applied magnetic fields, they emit electromagnetic radiation that is known as synchrotron radiation. This instrument is routinely used through-out the world and provides radiation in a wide spectral range from infrared to hard x-rays. Up to now, three generations of synchrotron radiation sources have been developed, with the fourth generation – called free electron lasers – already in use produces high power and ultrafast pulses of highly coherent radiation. Synchrotron based spectroscopy experiments have become indis-pensable to study the electronic structure of matter.

2.1.2 Storage Rings and Insertion Devices

Figure 2.2 shows the layout of a third generation synchrotron radiation fa-cility, the European Synchrotron Radiation Facility in Grenoble, France. Electrons are projected to relativistic speeds into a well-defined closed circu-lar path called storage ring that consists of straight and curved magnetic sections. The energy loss encountered by the electrons from emitted radia-tion is replenished by an oscillating electric field in a radio frequency cavity. The storage ring consists of alternating curved and long straight sections specifically optimized for producing intense radiation. Several different types of magnetic structures, called insertion devices, are tangentially placed to the storage ring to bend the electron beam and emit highly conical beams of light that are delivered to the beamlines.

Common types of insertion devices are the bending magnet and the undula-tor and depicted in Fig. 2.2(b) and 2.2(c) respectively. The bending magnet simply bends the electrons in a single curved trajectory, and a narrow cone of radiation with a continuous spectral distribution of is produced (e.g 100 – 1000 eV)62. In contrast, the undulator consists of a period array of magnets with alternating polarity of transversal or helical structure. The electrons oscillate in the periodic magnetic structure of the undulator by means of the Lorentz force and coherent x-ray pulses are emitted in a very narrow cone

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with small angular and spectral displacements. The spectral range from an undulator is concentrated into a series of discrete energies (undulator peaks) at which the intensity is greatly enhanced. The desired photon energy is cho-sen by varying the distance between the gap of the upper and lower magnetic arrays. In order to produce elliptical or circular polarized light, the shifting of magnetic rows in an undulator or a combination of two undulators can produce polarized synchrotron radiation.

2.1.3 Beamlines and Endstation

Once the coherent radiation is generated via a bending magnet or an undula-tor, the light is guided to the experiment in a segment called the beamline. A beamline is a series of precision optical elements that direct, monochroma-tize, and focus the divergent light beam onto the sample, all kept in ultra-high vacuum. The highly monochromatized beam is achieved by diffraction from crystals (hard x-rays) or reflection gratings (soft x-rays) through differ-ent slits and mirrors. The optical components are general to most beamlines, but also vary depending on the type of experiment conducted. The most common design in the soft x-ray range is using a spherical grazing mono-chromator (SGM) at grazing incidence. The divergent incoming beam is diffracted and focuses the desired photon energy onto the exit slit of the monochromator. The highly monochromatized beam is further focused with a refocusing mirror to guide the light onto the sample surface. For hard x-rays, monochromatization happens with high-quality Bragg crystals suited for a specific photon energy range.

Figure 2.2 (a) A birds-eye view of a synchrotron radaition ring with differentinsertion devices. (b). A bending manget . (c). A wiggler (All figures adaptedfrom Ref. [62])

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2.2 Principles of Photoemission Spectroscopy (PES)

2.2.1 General Formalism: The Photoemission Process

Figure 2.3 illustrates a schematic on the principle of the photoemission pro-cess, which is based on the photoelectric effect. When a photon with suffi-cient energy hv is illuminated on a solid, there is a probability that a bound electron is excited or released from the sample. The kinetic energy of the released photoelectron can be described as the difference between the initial ground state of the neutral atom and the final excited state of the ionized atom as:

where EK.E. is measured from the vacuum level, and is the work of func-tion of the sample, and EB.E. is the binding energy measured from the Fermi level (EF). Therefore, the characteristic binding energy of a core electron is determined by the incident photon minus the measured kinetic energy of the outgoing photoelectron. Since the core-level is characteristic to each ele-ment, XPS is useful for the elemental analysis in solids. Any deviation of the binding energy (B.E) of a core-electron from that of a free atom provides valuable information on chemical bonding of valence electrons. An electro-static hemispherical analyzer measures the kinetic energy, and the work function is set by the spectrometer that is electrically connected to the sam-ple. The energy of the incoming light ( ) is known from an x-ray source that is either determined by the anode material or from a tunable source such

Figure 2.3 Ionization process of a (a) core electron and (b) and valence electronat different incident energies. (c) Auger decay mechanism, and (d) radiativedecay with the emission of a photon after an absorption event.

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as a synchrotron. The core principle is the identification and speciation of each photoemission process that can be identified to a specific atom. After a core-hole is created, the excited state will decay rapidly by a transi-tion of an electron from higher lying shell to fill the core hole. This decay process, a recombination of an electron with a core hole can either be non-radiative, i.e. Auger electron decay (Fig. 2.3(c)), or radiatively by an x-ray photon (see Fig. 2.3(d)). Since the core-hole is often characteristic to a spe-cific element, the energy of the decay transitions (Auger electron or x-ray emission) can also be separated and identified. An Auger electron has a characteristic energy, distinguishable from a normal photoelectron, and is a commonly used spectroscopic tool. The transition probability between two states, one that is described in terms of transition from initial ground state to final state is given by Fermi’s Golden Rule:

The delta function (δ) ensures the conservation of energy. The perturbation operator H, the interaction Hamiltonian that causes the transition, is ex-pressed as:

where is the electron mass, is the electronic charge, and the speed of light. is the vector potential operation and is the scalar potential of the exciting electromagnetic field. is the momentum operator of the electron. This first order optical process, described by the first term and a first order perturbation is a suitable description of a photoemission process. The second term can be neglected as it describes a two-phonon process. During the photoemission process and the creation of a core hole, the sur-rounding electron cloud of the remaining ion contracts (relaxes) in order to screen the positive charge. The subsequent relaxation may lead to the so-called shake-up and/or shake-off process, which are also observed in the photoemission spectra.

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2.2.2 Electron Escape Depth and Photoionization

After the local absorption of a photon, the photoelectron must propagate through the sample to the surface. In route to the surface, the excited photoe-lectron suffers energy losses through interactions with other electrons or collective motions in the surrounding medium. The magnitude and energy dependence on the mean free path, , is dependent on the scattering mecha-nism in the samples. This important parameter reflects the mean distance between two inelastic collision of the photoelectron, which is a material and photon energy dependent parameter. The escape depth for all materials roughly follows a “universal curve” as a function of incident energy shown in Fig. 2.4. For soft x-rays, the photoelectron ranges will be in the order of 5-10 Å, which makes the soft x-ray range more surface sensitive, and hard x-rays more bulk sensitive.

2.2.3 Resonant Photoemission Spectroscopy

When the incident photon energy is tuned near the ionization energy of a core hole, the emission from this resonant state is substantially modulated. This enhancement of is due to an interference of the core-hole and direct valence band photoemission. This phenomena leads to a pronounced effects for strongly interacting electrons systems where there is a giant enhancement in the photoemission cross section at resonance. This feature is most signifi-cant in systems that contain partially occupied d shells that results in strong Coulombic interactions among valence electrons. For example, when the incident photon energy is equal to the energy difference between the p core

Figure 2.4 Inelastic mean free path of electrons as a function of electron kinetic energy.

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level and the valence d states, a direct photoemission process of a valence d electron occurs as:

The resonant process includes another intermediate step and is expressed as:

Both the initial and final states of two processes have the same electron con-figurations. A quantum-mechanical interference occurs between the valence photoemission and core-hole excitation with Auger decay and thus a sub-stantial change in measured intensity. The electron emission is via an Auger-type path, where a photon excites the system into an intermediate state with a core hole, and the core hole decays via Coulomb interactions in an Auger-like autoionization process. This leads to two new possibilities for the elec-tron spectra. One channel is characterized by involving the bound state pho-toelectron, thus called participator decay, and the other channel with an electron in the normally unoccupied levels in the final state with two elec-trons removed in an Auger-like transition termed spectator decay. An im-portant distinction is the constant kinetic energy for the participator decay, whereas spectator decay shift in binding energy as the photon energy is scanned across an absorption edge. A schematic illustration of the resonant photoemission process via the absorption process is depicted in Fig. 2.5. Since the enhancement takes place only for the emission of d electrons, one can obtain the d partial spectral density weight of the valence band63. This technique is extensively used in this thesis, particularly for work on bulk and thin film oxides discussed in Chapter 3 and 4.

Figure 2.5 Schematic illustration of the resonant photoemission process, in which a core electron is excited to a resonant unoccupied state, and the subsequent decay results in a final state the same as direct photoemission.

Core Level

Ground State

OccupiedValence Level

X-ray Absorption ParticipatorDecay

SpectatorDecay

DirectPhotoemission

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2.3 Principles of X-Ray Absorption and Emission Spectroscopy

The absorption and emission of x-rays in condensed or soft matter is rich in phenomena and offer a breadth of quantitative analysis. The impact of XAS and XES on a variety of scientific disciplines has offered new insights that are supported by different levels of theory, much of it in the last few dec-ades. The significant accomplishments are possible with advancements in x-ray optics, detectors and analyzers, as well as the availability of high-brilliance beamlines. These observations have pushed the field of condensed and soft matter physics to new heights.

2.3.1 Introduction to XAS and HERFD-XAS

As with other inner-shell spectroscopies, x-ray absorption (XAS) involves the excitation of a core-level electron to unoccupied states of a solid as shown in Fig. 2.6(a). The energy of the incoming photon is of the same order as the binding energy of a core electron. As a result, the excited core elec-trons are a reflection of the mass absorption coefficient ( ) of a particular shell and an approximate symmetry-projected unoccupied density of states of the excited state, as shown in Fig. 2.6(b). The region before and after the absorption edge characterized by transition of the photoelectron to unoccu-pied bound states is called x-ray near edge structure (XANES). Furthermore, XANES absorption features are sensitive to local symmetry, oxidation state, and local site geometry of the transition metal ion. In the extended region of the absorption spectrum, a region fittingly named extended x-ray absorption find structure (EXAFS), the excited electron has significant kinetic energy and yields extensive quantitative analysis on the local geometry such as bond lengths, coordination, etc61.

Like XPS, XAS is a first-order optical process with its cross-section de-scribed by the Fermi Golden Rule stated in Equation 2.2. XAS spectra are usually measured indirectly by measuring the secondary process since it is much easier to perform experimentally. Highly excited core-hole states are relaxed by a variety of mechanisms, the two most significant being emission of an Auger electron and x-ray fluorescence. The absorption can be meas-ured indirectly by detecting the Auger electrons or, more commonly, by detecting electrons from multiple secondary processes called total electron yield (TEY). For conducting samples, this can also be done by measuring the drain current of the sample. The radiative relaxation is recorded in the form of a total fluorescence yield (TFY), the summed fluorescence yields of all dipole-allowed transitions to fill the core-hole. Much of the information ob-

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tained from TEY and TFY techniques yield the same absorption spectrum, however differ in their sensitivity to bulk and surface information. The es-cape depth of photoelectrons is rather shallow (~ 50 Å) and limits the infor-mation to the top-most layers of the surface, thus restricting TEY method for surface sensitive studies. Fluorescence emission comes from deep within the bulk of the samples, where the escape depth of the fluorescent photons is on the same order as the penetration depth, i.e. of ~ 1000 Å. Therefore, this technique is very suitable for bulk properties and buried layers. The major drawback of this technique is the saturation and self-absorption effects that deviate from a linear relation to the absorption coefficient. These effects are mitigated by changing the experimental geometry to grazing incidence as well as quantitate correction using reference compounds. Recently, an alter-native method was developed in order to measure an absorption spectrum based on the integrated x-ray emission from a different element in the sam-ple64, termed inverse partial fluorescence yield (IPFY). IPFY is linearly pro-portional to the x-ray absorption cross-section without any corrections due to saturation or self-absorption effects65.

There are two characteristic features in the edge region of a XANES spec-trum: 1) an intense absorption edge, and 2) a weak pre-edge feature seen just before the onset of the main absorption line as illustrated in Fig. 2.6(b). For 3d transition metals, the intense features from dipole allowed transition give information related to the oxidation state of the ion, whereas the pre-edge feature provide additional information on the symmetry of the system66. This region is particularly suitable for systems that have a loss of inversion sym-metry, such as ferroelectrics. Also the energy position, intensity, and the number of pre-edge peaks also provide information on oxidation state, cova-lency, and coordination number61.

A relative new technique that also indirectly measures the x-ray absorption of a particular shell with substantial improvement in resolution is high-energy fluorescence detected XAS, or HERFD-XAS67,68. In this scheme, one

Figure 2.6 (a) The absorption transition resulting in the excitation of a core elec-tron. (b) The different regions of the x-ray absorption spectrum that bear uniqueinformation as a function of incident energy. (c) The characteristic regions intypical x-ray absorption spectra.

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is able to detect emission form a single dominant decay channel, i.e. the Kα line of a selected element’s absorption edge, thus making this a partial fluo-rescence yield technique. The intensity variation of the tuned fluorescence line is scanned across an absorption edge. This secondary process detection is realized with a wavelength dispersive instrument to discriminate between various decay channels, e.g. only the Kα fluorescence lines. The data acquisi-tion principle is shown in Fig. 2.7 (left) for the example of a scanning Johann type spectrometer. This fluorescence detected XAS is a second order optical (photon-in/photon-out) process that is theoretically described by the Kra-mers-Heisenberg equation (see next section). The lifetime of the final state core-hole governs the overall resolution of the spectra, in the HERFD-XAS case, is the core-hole state of the decaying emission line. This detection scheme has the tremendous advantages offering high single-to-background round ratio and improved resolution. The broadening is reduced to the core-hole state of the recorded emission line rather than the core-hole of the excit-ed photoelectron, as demonstrated in the spectrum shape in Fig. 2.7 (right).

2.2.2 X-Ray Emission Spectroscopy

Following an absorption event, the excited atom is left with a core-hole that relaxes through radiative (photons) and non-radiative channels (electrons). The emitted energy of the x-ray photon is the difference the energized core-state (Ei) and a higher final state (Ef). The intermediate state of non-resonant x-ray emission exhibits a continuum electron, in the photoelectron state | ⟩ with energy ϵ. The recorded fluorescence spans a broad energy range and dominated by the dipole selection rules: and and expressed by the Kramers-Heisenberg relation69 (Equation 2.4). This formalism holds

Figure 2.7 – Illustration of an energy-dispersive spectrometer in a Johansson geometry used for HERFD-XAS technique (Figure adapted from Ref. [67]). (b) Comparison between total fluorescence yield (TFY, black) and high-energy resolution fluorescence detected x-ray absorption (HERFD-XAS) spectra of MnO at the Mn K absorption edge.

Inte

nsity

(arb

. uni

ts)

658065606540

Incident Energy (eV)

TFY HERFD

Mn(II)OMn K Edge

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if one assumes the incident photon energy Γ is well above the x-ray absorp-tion threshold. The intensity of the emitted energy covers a broad range of fluorescence lines and satellite states as shown in Fig. 2.8. The process there-fore shows little dependence on the energy that is used for the initial pho-toionization process. The subsequent decay directly maps the local occupied partial density of states due to the atomic character of the core levels and the dipole selection rules. This information is complementary to the unoccupied partial sates of an XAS spectrum, as both provide information on local elec-tronic structure and symmetry of the absorbing atom.

X-ray emission spectra are often divided into two main energy regimes thatrequire altogether different instrumentation: soft x-rays (<1500 eV) and hardx-rays (> 3 keV). They are also characterized between lower in energy core-to-core transitions (CtC-XES) and weak and sensitive valence-to-core transi-tion (VtC-XES), as show in Fig. 2.8 for a complete collection of a K edgeemission. The main transition fluorescence transitions, mainly Kα and Kβtransitions, shown on the lower energy side in Fig. 2.8, bears information onsymmetry, local spin density, and indirectly the valence shell electron con-figuration. The VtC peaks, exceedingly weak and higher in energy, probesthe valence levels of the system and provides complementary information tophotoemission spectroscopy due to the element-specific probe70.

In the present work, the absorption and emission of K and L edges are probed for TM cations and halide anions. The subsequent decay of fluores-cence lines from shallow levels and valence levels are monitored to interpret the partial valence states and local structure.

Figure 2.8 – X-ray emission spectrum marked with Kα and Kβ core-to-coreemission lines along with satellite Valence-to-Core (Kβ2,5 and Kβ’’) emissionlines (left). A schematic that depicts the transitions recorded in an XES spectrum(right).

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2.2.3 Resonant Inelastic X-Ray Scattering

RIXS is a photon-in/photon-out technique that maps the energy and momen-tum transferred to a system. The changes in energy, momentum, and polari-zation of the scattered photons are transferred to intrinsic excitations of the material under study, which are then detected via the RIXS process. Using conservation laws and selection rules, one can determine the dispersion rela-tions between separate excitations in the sample. Common to many transi-tion metal oxides and their archetypes, a range of excitations such as charge, orbital, spin, lattice, and even coupled ones are observed in the acquired spectra71–74. Much of the advancement in RIXS has benefitted from high-flux synchrotron sources and remarkable high-resolution instrumentation such as beamline monochromators and spectrometers.

In order to derive the basic formula describing the RIXS process, let us con-sider an x-ray scattering process that involves a photon with energy hνin in-cident on a material that then decays radiatively with photon energy hνout. The perturbation to the Hamiltonian of the photo-excited electron accounts for a first-order and second-order terms that define the scattering cross-section. This transition probability from the ground state to a final state in-duced by the two Hamiltonian terms, by virtue of the Fermi’s golden rule, produces a second-order process that requires an intermediate state before de-excitation. This combined excitation and de-excitation process that is coherently correlated by the Kramers-Heisenberg equation:

where and represent the initial and intermediate states with energies and respectively, while and correspond to the final state con-

figuration and energy. The transfer operator is T and broadening of the over-all spectra is denoted by , the intermediate state of the second-order pro-cess. RIXS implies choosing the excitation energy, usually close to a charac-teristic absorption edge of an element, where the cross section in Eq. 2.4 changes dramatically and gives rise to additional features than non-resonant XES. Several different channels can emerge at different incident x-ray ener-gies. Correspondingly, the overall RIXS spectra contain both signatures of photons emitted elastically, i.e. with the same energy as that of incident pho-tons, and with inelastic photons, for which the emitted x-ray energy has changed. Overall, the scattering process leaves the system in a charge-neutral but excited state.

38

Figure 2.9(a) visualizes the RIXS process by an energy level diagram that is conceptually separated into two steps. The first process involves an excita-tion of a core electron to an unoccupied band, and subsequent decay of an occupied valence electron. The indirect RIXS creates a scattering of the lo-calized core-hole wave of the photoelectron with the valence electrons, leav-ing behind a range of excitations in the valence electrons mentioned previ-ously. The features of these excitations take shape on an energy loss spec-trum that is a measure of energy transferred to the sample. The elastic line is the zero energy loss to system ( ) and is common to almost all RIXS spectra and varies in intensity at different absorption edges. The ine-lastic peaks are more interesting, as they provide clues into the interactions between the different quantum parameters of the system.

The characteristic traits of RIXS features are highlighted in this section. In general, the entirety range of features – both elastic and inelastic is strongly dependent on the photon energy, geometry of the experiment, and the mate-rial system. The RIXS experiments performed in this thesis are on transition metal oxides at Ti K, Mn K, Fe L, and O K edges, and at I L1 edge for the halide perovskite compounds, covering both soft (100 – 2000 eV) and hard x-ray (E > 2000 eV) edges. The separate excitations that are probed in thisthesis are charge transfer (CT) and d-d excitations. Charge-transfer excita-tion covers an energy range between 4-10 eV in energy loss ( )that are typically observed for transition metal compounds due to the broadband hybridization between the metal d and ligand p states. Assuming thatthe system is excited to a condition, this results in a intermedi-ate state where is the core-hole in the ligand. This feature is obtained whenthe de-excitation process leaves behind an electron the TM d orbital with ahole in the valance band, and reflects the O 2p – M 3d hybridization.

Figure 2.9 (a) Schematic view of the RIXS process. (b) Illustration of a 2p → 3d RIXS process that probes two finals states: a charge-transfer from O 2p → TM 3d and/or excitations within the d manifold.

39

The d orbitals are directly probed at the L edge in RIXS experiments. The d-d excitation is in the range of 1-3 eV energy loss which correspond to elec-tron-hole transition between the ligand-field non-degenerate d states. This transition is generally forbidden due to the dipole selection rules, however are observed due to the coherent second-order nature of the RIXS process. Once a d orbital is reached from an excitation of the L edge, the intermediate state is an excited electron that is strongly bound to the core hole that does not easily delocalize to neighboring sites. The subsequent de-excitation is the decays from a different d orbital filling the core-hole so that the number of electrons in the system is preserved. This excitation is especially strong in transition-metal oxides which manifest strong correlations. The excitation process leaves a change in the occupancy of the d orbitals that is largely determined by the crystal field splitting.

2.4 Thin Film Fabrication of Ferroelectrics

2.4.1 Pulsed Laser Deposition

Perovskite oxides have a simple crystalline structure that permits the epitaxi-al growth onto single crystalline substrates. The primary technique to grow oxide thin films is pulsed laser deposition (PLD) technique. This technique allows fabrication of thin films using stoichiometric targets of the same compound for growth and is relatively simple from the application context. In reality, the interaction of the laser-matter interaction and the subsequent generation of plasma is dynamic and very complex and not entirely under-stood75.

Figure 2.10 A schematic illustration of the PLD deposition process(left). A plasma plume produced from a BaTiO3 target immediately afterablated by a 248 nm KrF laser (right).

40

PLD is a physical deposition technique in which an intense laser ablates and vaporizes the target material to form onto a single crystalline substrate. The stoichiometric transfer is ensured as long as the ablation is homogenous. The laser used for this thesis is a UV high-power pulsed laser based on an excim-er KrF (248 nm). The principles of operation are: 1) the incoming laser va-porizes the target material, 2) the generation of a plasma that expands rapidly and dynamically interacts with the background gas, and 3) nucleation of a solid phase from a supersaturated gas that reproduce the target composition onto a substrate. The desired nucleation of the material onto the substrate can be adjusted by a number of independent parameters. These include the tar-get-substrate distance, the laser energy density and repetition rate, the oxy-gen partial pressure, and the temperature of the substrate. The nucleation of the target material constituents on a sample surface determines the growth mode of the film: island growth (Volmer-Weber), layer-by-later (Frank-Van der Merwe), or Stranski-Krastanov growth76. Results on thin film deposition are reported in Chapter 4 of this thesis.

2.4.2 X-ray Diffraction and Reflectivity

X-ray diffraction experiments were performed on a PanAlytical X’Pert four axes diffractometer. This setup is equipped with a Cu Kα x-ray tube with a wavelength of 1.5406 Å and x-rays emitted onto the sample with a parallel beam configuration, both for x-ray reflectivity and out-of-plane diffraction measurements. The two separate measurements were used to obtain out-of-plane lattice parameters, understanding of the crystalline and epitaxial quali-ty, and thickness of the films. For x-ray diffraction measurements, a PiXel detector was used and a proportional detector consisting of a xenon/methane mixture for XRR measurements.

2.4.3 Piezoforce Microscopy

Scanning probe microscopy is used to provide information on the surface morphology of deposited thin films. Atomic force microscopy (AFM) and the more advanced piezoresponse atomic force microscopy (PFM) have be-come reliable tools to study ferroelectric thin films77. In the AFM technique, a small tip is brought to the proximity of the film surface and the defor-mation of sample surface is monitored. A collimated laser beam reflects the cantilever of the AFM probe and is reflected to a photodetector to image the sample surface. The image can provide clues to the sample roughness, ho-mogeneity, and type of film growth.

41

Piezoresponse is an extended type of force microscopy that measures the dynamic electromechanical response of a ferroelectric sample. Ferroelec-tricity is commonly demonstrated by the PFM hysteresis loops and a phase difference of the acquired images before and after poling78–80. The technique involves the application of an AC voltage to the conducting tip in mechani-cal contact with the surface and the subsequent oscillation of the tip is de-tected by a lock-in technique. Due to the inverse piezoelectric effect, the ferroelectric material responds with an expansion or contraction, which is measured as the cantilever deflects through a laser spot. During PFM meas-urements, the measured amplitude and phase of the deflected signal contains information about the local ferroelectric strength and the orientation of the ferroelectric domain. The PFM signal yields local hysteresis loops and the switching voltage and can be extracted. Results on the ferroelectric properties of BTO and doped BTO thin films are presented in Chapter 5.

42

3 Bulk Perovskite Oxides

In this section we review the work done on two bulk perovskite oxides: Mo doped LaFeO3 and Sr2FeMoO6. The electron doped LaFe1-xMoxO3 (LFMO) exhibits a novel ground state that combines room temperature metallicity with antiferromagnetic ordering81. The second compound, Sr2FeMoO6

(SFMO) has received considerable attention since the discovery of a tunnel-ing magnetoresistance (MR) and spin-polarized metallicity at room tempera-ture82. A combination of x-ray spectroscopies is used to study the valence electronic structure. Both these materials offer an attractive possibility for their use in magnetic devices and spintronics83,84, which is the intention of our future work. The motivation for this work here is to comprehend the electronic structure through different x-ray techniques and reveal the conditions that make two materials appealing for spintronic applications.

3.1 Insulator-to-Metallic Antiferromagnetic Oxide: LaFe1-xMoxO3

In Paper I, we report on a new material that exhibits a rare ground state in a perovskite oxide. By doping at the B-site of the LaFeO3 oxide with the 4d Mo transition metal (LaFe1-xMoxO3), the system undergoes a gradual crosso-ver from an insulating to metallic state at room temperature while retaining its antiferromagnetic (AFM) spin ordering. This new compound was investi-gated through several laboratory and synchrotron-based techniques to show its unique ground state properties. This material may hold promise in its use in antiferromagnetic spintronics, as well as fundamental questions in regards to its unusual coexistence of metallicity and AFM at room temperature. Six samples of LaFe1-xMoxO3 (LFMO) with Mo doping at x = 0.00, 0.05, 0.10, 0.15, 0.20, 0.25 were synthesized by the solid-state route. Magnetic measurements were carried out using a Quantum Design SQUID magnetom-eter, as shown in Fig. 3.1(a). The AFM transition Neel temperature (TN) for the doped compounds decrease linearly with increased Mo doping (see inset of Fig. 3.1(a)). This indicates that even at high doping concentration, the

43

AFM order is still maintained. Changes in resistivity should also accompany this gradual metallic transition and has been measured using a four-point probe setup as shown in Fig. 3.1(b). There is a substantial decrease in the resistivity magnitude for LFMO-0.25 compared to LFMO-0.10, which re-mains insulating. Although the slope negative slope suggests an insulating behavior (a negative ∂ρ/∂t), it does not diverge and saturates at low tempera-tures. This resistivity is attributed to the formation of an insulating Mo6+ that exists at the grain boundaries, as shown by several previous studies and is a common observation in polycrystalline ceramics82,85. This is further investi-gated through scanning tunneling microscopy (STS) measurements shown in Fig 3.1(c). A finite dI/dV at zero bias (V = 0) exists for both 20% and 25% Mo doping, indicating a finite DOS at EF for both samples, which is insignif-icant for x = 0.0, 0.10, and 0.15 at V = 0. The effect of surface treatment is resolved for the metallic LFMO-0.20 sample, where the untreated sample revealed negligible dI/dV intensity (see inset Fig. 3.1(c)) compared against the in-situ cleaned LFMO-0.20 sample. This data set corroborates with the measured photoemission spectra and confirms the crossover to metallic state through electron doping.

The mechanism proposed for the electronic insulator-to-metallic transition can be contended to the mixed-valency in the Fe ions and concurrent struc-tural distortions. In the pure LaFeO3, the ground state ionic picture for Fe is a Fe3+ ion that is localized and positioned equidistantly from its’ six adjacent oxygen atoms. The mixed valency is probed through Fe L x-ray absorption spectroscopy shown in Fig. 3.2(a) for all samples. As evidenced from the fitted data for x = 0.20 spectra in Fig. 3.2(b), the lower energy peak is pre-dominantly Fe2+ while the peak at higher photon energy is attributed to Fe3+. The Fe2+ component monotonically accrues as the Mo doping concentration steadily increases.

Figure 3.1 (a) Temperature dependence on magnetization for all compositions.(b) The resistivity vs. temperature plots and (c) Scanning tunneling spectroscopy(STS) plots on the Mo doped samples. The inset compares the STS for an un-treated x = 0.20 sample.

44

To assess the changes to Mo valency, PES was used to probe the Mo 3d core level, as shown for LFMO-0.15 in Fig 3.2(c). The spectrum is fitted consid-ering different Mo oxidation states and a satisfactory fit only with three Mo components (Mo4+, Mo5+, and Mo6+) to match the experimental spectra. The average valency for both Fe and Mo ions is extracted for all samples using these two core-level spectra. The obtained numbers of doped electrons at the Fe and Mo sites relative to their most stable valency of 3+ and 6+ respective-ly, are plotted in Fig. 3.2(d). With increasing doping concentration, electrons gradually fill both Fe and Mo bands. However, the electron doping rate at the Fe sites reduce after x = 0.15 and saturates to ~0.5e-/f.u. for x ≥ 0.20. Further inclusion of Mo ions to the system, i.e. as for the case of the x = 0.25 sample, pushes all additional charge to the Mo sites, as evident from the identical spectra for x = 0.20 and x = 0.25 in Fe 2p XAS in Fig. 3.2(a). These observations imply that the electron doping at the Fe sites reaches a thresh-old when the valency becomes 2.5+ exceeding 20% Mo doping. It can be implied that the 2.5+ average valency forces an equal sharing of one electron between two Fe-sites. From our theoretical work (see Fig. 3.2(e) - (g)), the Fe states are found to be insulating and localized, which also suggest an equal number of Fe3+ and Fe2+ ions in the system for x = 0.25. This provides further evidence for a charge disproportionation in this system where itiner-ant and localized orbitals reside on different atomic sties.

Figure 3.2 (a) Fe L edge XAS spectra for all compositions that are in-situscraped with a diamond file. (b) Fitted data for x = 0.20 sample by the weighted average in terms of Fe2+ and Fe3+. (c) Mo 3d core level XPS spectra for x = 0.15. (d) Different contributions of doped electrons obtained at Fe band, Moband, and the total/f.u. and nominal/f.u. (e)-(g) Optimized GGA+U total andpartial density of states showing the two charge disproportionated Fe sites.

45

3.2 Valence Electronic Structure of LaFe1-xMoxO3

The results on the valence electronic structure for Mo-doped LaFe1-xMoxO3 are taken from a working manuscript (Paper II). The purpose of this study is to combine several x-ray spectroscopies to resolve the ground state electron-ic structure and connect the concomitant structural distortions as a function of doping.

Figure 3.3 shows the experimental valence band as measured by photoelec-tron spectroscopy measured at an incident energy of 2500 eV for LaFe1-

xMoxO3 compounds (x = 0.0, 0.10, and 0.25). A closer inspection of the va-lence band near the Fermi level (EF) shows an increase in spectral weight as Mo content is increased. For the parent compound LaFeO3, a sizable gap between the valence band maximum (VBM) and the Fermi level (EF), in agreement with a previous soft x-ray study on LaFeO3

86,87. For x = 0.10 doped sample, a nominal increase in states appear at 1 eV in binding energy (B.E.), however it is negligible at 0 eV, consistent with its insulating proper-ties. Conversely, for the x = 0.25 sample, we observe a sizeable intensity across EF, confirming the transition from an insulating LaFeO3 to a metallic state by sufficient doping at the B site81.

A more comprehensive study on the valence electronic structure has been probed through resonant photoelectron spectroscopy at the Mo L3 and Fe L3 edges. The element-selective projection to the partial density of states (DOS) through RPES spectra allow for identification to which TM states constitute the valence band and the extent of their energy positions and hybridization.

Figure 3.3 HAXPES data measured at 2500 eV on the LaFe1-xMoxO3 series. Theinsulating LaFeO3 and x = 0.10 compounds exhibits a gap between the valenceband maximum and the Fermi level (EF), whereas for the x = 0.25 composition,there is clear density of states at EF, characteristic of a metal.

46

Fig. 3.4(b) shows RPES spectrum collected at range of incident energies across the Mo L3 absorption edge (Fig. 3.4(a)). Three structures are enhanced when the energy approaches the absorption maximum. The spectral region around 0 eV B.E. shows an intense and narrow peak (peak labeled A) that can be assigned to Mo t2g states81, where the itinerant electrons in the partial-ly filled sub-band gives rise to the metallicity for x = 0.25. The broad en-hancement in the 5-8 eV range signals Mo 4d - O 2p admixture as was shown by our band structure calculations in the previous study81. These spec-tra qualitatively show the emergence of new Mo 4d states that establish the metallicity in this material for doping levels at x = 0.25.

The RPES spectrum collected across the Fe L3 absorption edge is shown in Fig. 3.5. The resonant spectra show four principal structures that are en-hanced as the excitation energy is tuned to the resonance condition (709. 0 eV). The intensity around EF shows no enhancements at the absorption max-imum, indicating that the states at EF have no contributions from Fe, but are solely due to Mo 4d states, as discussed in the previous section and corrobo-rated by band structure calculations81. The inset in Fig 3.5(b) shows a narrow range of the valence band expanded at a few selected energies (703.0, 707.5,

Figure 3.4 (a) Mo L3 absorption spectrum for LFMO-0.25 measured in TEYmode. The tick marks indicate the excitation energies used to collect photoemis-sion spectra. (b) RPES spectra for LaFe0.75Mo0.25O3 showing three enhancedspectral features that relate to Mo 4d contribution in the valence band.

47

709.0, and well above resonance at 712.1 eV). When the incident energy is tuned to the maximum of the shoulder peak, associated with Fe2+ species, an enhancement for states just below the EF is observed (see peak D). Con-versely, at the absorption maximum spectrum (709.0 eV), the structure E near 3 eV in binding energy gains substantial intensity in contrast to the off-resonance energy (703.3 eV) - with much more than at excitation energy corresponding to the peak representative of Fe2+. These two structures may point to spectral weight emerging from the two inequivalent Fe sites, i.e. the local charge disproportionation of Fe in the metallic compound.

RIXS spectrum was acquired for incident energies hνi across the Fe L edge for three compounds in the series and shown in Fig. 3.6. There are several well-resolved structures that can be readily identified based on the energy loss of the excitations involved. At low energy losses, both the two insulat-ing LaFeO3 (Fig. 3.6(a)) and x = 0.10 (Fig. 3.6(b)) doped samples exhibit three constant energy loss (also referred to as Raman-like) features at -5, -3.1 eV, and-1.8 eV (indicated by dashed lines) that signal intra-site Fe d-d exci-tations. These are typical of strongly correlated systems as observed in many other iron-based oxides88–90. The fixed energy loss features are assigned to

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Figure 3.5 (a) The Fe L2,3 XAS spectrum for x = 0.25 with a strong shouldercentered at ~ 707.5 eV. (b) RPES data collected at several incident energies withfour principal structures. The inset shows features that are the directly related tothe charge disproportionate Fe sites.

48

electronic excitations which involve localized states while excitations in-volving delocalized character linearly disperse with incident energy (hνi)71,72. RIXS on LaFe0.75Mo0.25O3 (Fig. 3.6(c)) clearly shows strong fluorescence contributions at low energy losses that are superimposed on top of the d-d excitations, which is a characteristic feature of metallic sample91, where low energy excitations across the Fermi level are possible. As the fluorescence disperses to higher-energy with increasing incident energy, well-resolved features appear in its low-energy loss tail, with distinct d-d features at -3 eV, -1.6 eV, and -0.8 eV and remain at fixed energy loss with increase in inci-dent energy. The peak positions shift to lower energy in metallic LaFe0.75Mo0.25O3 compared to that of insulating LaFe0.90Mo0.10O3 and La-FeO3. The CT features are distributed across the different energy losses en-ergies above -6 eV; this is in accordance with the modified Fe-O hybridiza-tion for the shorter Fe-O bond distance81. Furthermore, the RIXS spectra for the metallic sample exhibit clear low-energy loss intensities below 1 eV, providing a fingerprint to closing of the insulating gap. This signature has also been observed in many insulating-to-metallic transitions such as in NdNiO3

91 and La1-xSrxMnO392.

The comprehensive x-ray spectroscopy results reveal an unusual metallic state with itinerant (Mo d) and localized carriers (Fe d) in LaFe1-xMoxO3.

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Figure 3.6 Fe L edge RIXS spectra for (a) the undoped parent LaFeO3, (b) x = 0.10 Mo doped sample, and (c) the metallic crossover at x = 0.25 at selected photon ener-gies. The vertical dashed lines mark energy positions of inelastic loss features and arrows are used to guide the linearly dispersing fluorescence-like structure.

49

3.3 The Double Perovskite Sr2FeMoO6

The results discussed in this section are obtained from x-ray spectroscopy measurements taken from Paper III. Sr2FeMoO6 (SFMO) has gained note-worthy attention due to its large changes in resistivity through application of small external fields – at room temperature82. Its electronic and magnetic properties are directly linked to the two TM cations and their interaction with each other and adjacent oxygen atoms93. Significant experimental and theoretical work has been done with the aim of describing the complex mi-croscopic origin of its physical properties, however uncertainty still remains in regards to its ground state ionic configuration.

In a completely ordered SFMO structure, the Mo and Fe alternate in a rock salt type order that repeats in all crystallographic axis as shown in Fig. 3.7(a). The magnetic ordering is usually considered as a ferrimagnetic state with parallel Fe moments that are antiferromagnetically aligned to Mo mo-ments. This material is also noteworthy for its half-metallicity, meaning the density of the states for the down-spin band is present at the Fermi level, whereas the up-spin band forms a gap as shown in the schematic in Fig. 3.7(b). In order to disentangle the Fe 3d and Mo 4d states in the valence band, reso-nant photoemission is aptly suited to separate the orbital-resolved d contribu-tions. Figure 3.8(a) shows the RPES spectra taken across the Mo L3 absorp-tion edge measured in TEY mode. The L3 structure is formed by a doublet structure at 2521.6 eV and 2525 eV (Fig. 3.8(a)), which agrees well with a previous study with multiplet calculations that assume a Mo5+ (4d1) valency94. In the RPES spectra shown in Fig. 3.8(b), one can easily distin-

Figure 3.7 Schematic of the tetragonal crystal structure for SFMO with alternatingFeO6 and MoO6 octahedra in an ordered perovskite. (b) Schematic electronic struc-ture that show metallicity from spin-polarized carriers in the spin-down channel.

50

guish three well-defined structures labeled A, B, and C that are all modified by the incident energy across the Mo L3 absorption edge. A finite density of states appears between 0.5 eV and the EF (feature labeled ‘A’) at all photon energies, and enhanced at the maximum of the absorption threshold (labeled ‘8’ in top panel of Fig. 3.8(a)). This prominent feature relates directly to itinerant Mo d states at the EF, in agreement with previous results95. The broad B feature ranging from 0.8 – 2 eV can be assigned to an admixture of spin-up Fe eg↑ states and spin-down Mo t2g↓ states. The structure at 7-8 eV has also been argued to belong to Mo 4d states hybridized with O 2p bond-ing states96. These features together qualitatively show the energy positions and spectral density of the delocalized Mo d states in the valence band, a crucial component of the electronic description of SFMO. The Fe L absorption edge and the resonant photoemission spectra are both shown in Fig. 3.8(c,d). The off-resonance case can be regarded with energy of 704.5 eV, below the Fe L3 absorption threshold. The spectrum clearly shows a tangible number of states at the Fermi level due to Fe d states (Fea-ture ‘A’), which can be assigned to the delocalized Fe t2g↓ spin states. The shoulder feature B, separated by about 1.25 eV from EF, is attributed to the Fe eg↑ spin-up states. The energy separation between feature A and B can be associated with the spin-down t2g↓states and spin-up eg↑ states respectively. The significant enhancement of the states between 3 – 10 eV (labeled C and D in Fig. 3.8(d)) are due to strong Fe 3d – O 2p hybridization. Feature C and

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Figure 3.8 (a) Mo L3 XAS in TEY mode. (b) Mo RPES reveal three structures(marked ‘A’, ‘B’, and ‘C’) that related to Mo d states in the valence band. The inset shows a narrow VB range at off-resonance on resonance photon energies. (c) Fe L3 XAS spectrum. (d) Fe RPES show four features (labeled A-D) that are enhanced atthe absorption maximum (~ 708.9 eV) in contrast to incident energy below the absorption threshold (704.5 eV). The A and B features can be assigned to the t2g↓ and eg↑ respectively.

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D at 5.5 eV and 4 eV respectively are ascribed to Fe t2g↑ and the Fe eg ↑ bonding states that hybridize with the O 2 p states. These spectral signatures fall exactly in line with previous absorption and photoemission experiments95, and are in line with proposed models that explain the metal-licity in SFMO82,93. Accordingly, the RPES data shows that states with in-tensity at the Fermi level are due to both Fe 3d and Mo 4d states.

RIXS is another technique that can directly probe the valence electronic structure as well as the correlated nature of d orbitals. Fig 3.9(a) shows the near-edge XAS for SFMO and two related oxides with a Fe3+ valence struc-ture measured in both TEY and TFY. All three samples show a low-energy shoulder, which is common to compounds with a Fe3+ formal valence 97,98. The two groups of peaks near 710 eV (L3) and 721 eV (L2) are due to spin-orbit Fe 2p core levels, which are further split by the interplay of crystal-field splitting, charge transfer, and p-d/d-d interactions multiplet effects 99,100. The RIXS spectrum is collected across the Fe L absorption edge is presented in a 3D waterfall plot in Fig. 3.9(b). The RIXS map is useful to visualize the low-energy excitations of the system as a function of different incident ener-gies. The RIXS map in Fig. 3.9(b) show a constant emission energy fluores-cence features centered at ≈ 709 eV and near ≈ 720 eV that are Lα and Lβ lines respectively. To further analyze the RIXS spectra of SFMO, we compare the different energy loss features against the insulator α-Fe2O3 that also displays a weak magnetism at room temperature101, as shown in Fig. 3.10. SFMO exhibits

Figure 3.9 (a) XAS measured in TEY and TFY for SFMO and two reference Fe3+ compounds LaFeO3 and α-Fe2O3. (b) A RIXS waterfall plot measured on an emis-sion energy axis.

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very broad features clearly contrasting the well-resolved peaks for the insu-lating α-Fe2O3. The two Raman-like energy loss features centered at approx-imately 1.4 eV and 3.0 eV (marked α and β respectively in Fig. 3.10) for α-Fe2O3 can be assigned to intra-band d-d transitions between the Fe3+ t2g and eg bands. The same features are also distinguishable for SFMO, although broader than that of α-Fe2O3. The inelastic features for SFMO are superim-posed on top of fluorescence emission. A closer inspection reveals that the peaks that are well defined for α-Fe2O3 can also be tracked at constant ener-gy loss with increasing incident energies for SFMO, particularly at excitation energy of 716.2 eV. These features are typical for crystal field excitation in correlated materials with 3d transition metal elements71,72. The early onset of the fluorescence-like structure on the SFMO sample can be attributed to the half-metallic nature of the sample. The larger bandwidth of the inelastic structures for SFMO in the low energy excitation range, from 3 - 6 eV, must also include inter-band transition from Mo d manifold to Fe t2g and eg states. This can be sensibly argued due to the extensive hybridization of Mo d states that appear alongside Fe d states as highlighted through RPES data. Despite these assessments, further studies of the Mo L RIXS, corroborated by ab initio calculations can provide key information on the coupling between Fe and Mo ions.

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53

4 Bulk and Thin Films of a Doped BaTiO3

In this chapter, a comprehensive x-ray spectroscopy results are presented on a co-doped ferroelectric BaTiO3 (BTO). Much of the work presented in this section is taken from the conclusions in Paper IV and V.

4.1 Low Bandgap and Ferroelectricity of co-doped BaTiO3

This chapter presents the work on bulk and thin film ferroelectrics based on BaTiO3. BaTiO3 (BTO) is one of the most well studied ferroelectrics and used in many different applications102–106. However, the large bandgap of BTO (Eg: 3.2 eV) can harvest only a tenth of the solar irradiation, making it unsuitable for the light absorption in the visible spectrum.

In this context, our research team used a double-element doping strategy to reduce the bandgap of BTO to the visible region of the spectrum without compromising the polarization. The co-doping at the Ti-site of BaTiO3 with trivalent Mn3+ and a d0 Nb5+ ions (BaTi1-x(Mn1/2Nb1/2)xO3) showed an appre-

Figure 4.1 Change in optical band gap with increase in concentration of dopants.Band gaps of a few standard photovoltaic materials are also indicated. Figureadapted from Ref. [107].

54

ciable reduction in the bandgap and a sizable retention of the remnant polari-zation107. The partially occupied Mn 3d4 should meet the prerequisite for bandgap variation30,108 and the charge-compensator Nb5+ may provide an off-centering similar to that in KNbO3

109. In Fig. 4.1, bandgap values for differ-ent compositions for BaTi1-x(Mn1/2Nb1/2)xO3 series (hereafter referred as BTMNO) is obtained through UV-Vis diffuse reflectance spectra. The bandgap continuously decreases from 3.18 eV for pure BTO to 1.3 eV for the composition with x = 0.15. The polarization was found to be largely pre-served, with a sizable room-temperature polarization of 15.4 μC/cm2 for the x = 0.075 sample, 70% of the parent BaTiO3 (25 μC/cm2). Here, we probe the changes in the electronic structure as a function of dop-ing using combined x-ray spectroscopies. First, we monitor the displacement of Ti atom using the pre-edge structure at the Ti K edge as shown in Fig. 4.2(a). The first two peaks in the magnified pre-edge spectra (A1 and A2 in Fig 4.2(b)) directly relate to the quadrupolar 1s → 3d t2g and mixed Ti eg - O 2p transitions respectively110,111,66. The intensity of peaks A2-A5 decrease progressively upon increased doping with Mn/Nb ions when compared to BaTiO3. The extent of off-centering of Ti atom can be directly related to the relative change in intensity of the pre-edge A2 peak112,113. The area under peak A2 is estimated with a combination of Gaussian and an arctangent background, exhibits the same trend as the PR as a function of doping as shown in Fig. 4.2(c). The increase in PR for the x = 0.075 sample compared to x = 0.05 is also reflected by a corresponding increase in the A2 peak area from the Ti K pre-edge spectra. This finding indicates that the Ti atom is still off-centered with respect to the octahedral cage of surrounding oxygen at-oms, thereby retaining ferroelectricity in presence of doping of 3d (Mn) and 4d (Nb) transition metals.

To understand the changes in the valence electronic structure as result of dopants, we use resonant photoemission at the Mn L edge. The resonant

Figure 4.2 (a) Ti K edge HERFD-XAS for all compounds. (b) Pre-edge feature thatshow a minor change in the intensity of the quadrupolar transitions (peaks A1 and A2) and indicate that the Ti off-centering is maintained. (c) The polarization behav-ior is consistent with measured A2 peak intensity for the bulk samples.

55

process at the Mn core 2p states enhances the Mn 3d spectral weight, making it feasible to characterize the Mn 3d states and their position in the valence band. Figure 4.3(a) shows the RPES on thin films collected at the maximum of Mn L2,3 absorption edge. As shown by the resonant enhancement, elec-tronic states near EF has Mn 3d character. The on-off resonance difference spectra, denoted by the gray shaded area (Fig. 4.3(a)), provides a direct probe for Mn 3d states. Although the spectral shapes are similar, the intensi-ties of all three peaks, C1, C2, and C3 (Fig. 4.3(b)) show a larger enhance-ment for x = 0.075 when compared with x = 0.025 film, i.e. increasing inten-sity with an increasing Mn content, as expected for these newly created Mn states. From the RPES data presented here, we find the dopants shift the valence band edge by as much as 1 eV relative to the Fermi level (EF). Thus, the growth of new states with predominantly Mn 3d character within the bandgap of BTO (marked C1 in Fig 4.3(b)) can be linked with the observed bandgap reduction. This finding corroborates with DFT+U calculations that clearly show formation of new states at the top of the valence band edge with predominant Mn 3d character (see Fig. 4.3(b)).

4.2 Thin Film Properties of co-doped BaTiO3

In order to realize the unique properties of BTMNO for use in ferroelectric photovoltaics, high-quality thin films are fabricated by pulsed laser deposi-tion (PLD). We explore Mn/Nb doping levels for x = 0.025, which in our previous study of bulk materials, was found to exhibit the highest saturation

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Figure 4.3 (a) RPES spectra for doped and undoped samples at the Mn L edge in the off-resonance (hνi = 630 eV) and on-resonance condition (hνi = 640.2 eV). (b) The three resonating features for x = 0.025 and x = 0.075 are assigned to Mn d states. (c) Total and partial orbital-projected density of states for BaTiO3 (top) andx = 0.11(bottom).

56

polarization of 22.3 μC/cm2, and x = 0.075 (lowest band gap) in addition to the parent BaTiO3 thin films. The thin films were fabricated on (001) SrTiO3 substrates with a 0.1˚ miscut to assure a layer-by-layer growth. The high-resolution diffraction pattern is shown in Fig. 4.4(a). The three clear diffrac-tion peaks correspond to (00l) reflections signifying epitaxial growth and no indication of any secondary phases is observed. The position of the (002) peaks systematically shifts toward lower angles with increasing Mn/Nb con-tent, suggesting the two transition metals occupy the B-site. The out-of-plane lattice parameters calculated from the (002) reflections are 4.366 Å, 4.376 Å, and 4.386 Å for BTO, BTMNO-25, and BTMNO-75 respectively. The a-axis lattice parameter for bulk BTO is 3.992 Å114, indicating that all three films are strained due to the mismatch with the STO substrate (a-axis 3.905 Å). The inset of Fig. 4.4(a) shows expanded region around the (001) peak, with reflections clearly displaying interference fringes suggesting a relative-ly coherent and a smooth interface with the SrTiO3(001) substrate. From the cross-sectional SEM images in Fig. 4.4(b), it can be observed that the thick-ness of the BTO layer is ~ 75 nm. AFM images taken across the film surface are shown in Fig. 4.5(a)(b). Since the samples are grown on an STO substrate with a 0.1˚ miscut, the step-terrace structure of the deposited film following post-annealing is clearly observed in the AFM image shown in Fig. 4.5(a). This clearly indicates atomically flat terraces with single unit-cell steps (~ 4 Å). For thicker films, 3D clusters nucleate on these layers due to the stress induced by the lattice mismatch (Fig 4.5(b))115. The bandgap of the films are characterized using a UV-Vis spectrometer in the total reflectance mode as shown in Fig. 4.5(c). The bandgap of BaTiO3 has previously been reported to be nearly 3.18 eV, just outside of the visible spectrum107. Here, the optical band gaps can be estimated from the tangent

Figure 4.4 (a) XRD scans of BTO and Mn1/2Nb1/2 doped BTO thin films. (b)cross-sectional SEM image showing the thickness and a uniform BTO film on SrTiO3.

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lines in the (αhν)2 versus hν plots for the three samples. The BTMNO-75 shows a bandgap of approximately 2.6 eV, enabling the absorption of sun-light in the visible range, lower than that for BTO (3.2 eV) thin film. This is in line with our report on polycrystalline samples, where the bandgap reduc-tion was shown to result from Mn d states that hybridize with O 2p states at the edge of the valence and conduction bands116. The ferroelectric hysteresis loops can be revealed through the PFM response. This technique has developed into a successful tool for direct evidence of ferroelcitrcity77,117,118. Figure 4.6(a)-(c) show a hysteretic behavior typical of polarization switching for all three samples. The top panel shows the ampli-tude switching and bottom panel shows the phase switching as function of bias. The PFM phase-voltage curve shows a fully saturated hysteresis for the

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Figure 4.5 AFM images of (a) 75 nm doped BTMNO-75 film with well-definedterraces indicating layer-by-layer growth. (b) 125 nm BaTiO3 thin films with 3Dnucleated clusters. (c) UV-Vis measurements in reflectance mode for BTO,BTMNO-25, and BTMNO-75 thin films with a bandgap tuning from 3.2 eV to2.6 eV.

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BTO and doped BTMNO samples. All curves show an 180˚ phase change across the saturation voltages, signifying complete switching, which is char-acteristic to a true ferroelectric. The hysteresis for BTMNO-075 shows a slight anti-symmetric phase with respect to negative DC bias, which can likely be attributed to defect-induced polarization pinning in the upward direction119–122. We note here that the hysteresis loop is offset towards a neg-ative voltage indicating a built-in electric field that is also found in similar studies119. The amplitude curves (Fig. 4.6 top panel) show a nearly symmet-rical butterfly shape for the BaTiO3 sample, which indicates a robust and switchable ferroelectricity. The curve is slightly anti-symmetric for BTM-NO-25 and BTMNO-75 thin film samples likely due the strong poling for doped thin films. Both hysteresis and butterfly loops however are character-istic of classical ferroelectric switching. To demonstrate the effect of poling on ferroelectric switching of the thin film, we reverse the upward polarization by scanning a diagonal line across the film surface with a tip biased with a voltage VTip = +5 V that exceeds the coercive voltage for the film. Figure 4.7 shows the resulting polarization pattern as measured by PFM for the BTMNO-25 sample. The antiparallel domain line shows a clear contrast with the polarization in the opposite di-rection, showing the polarization response is strong with domain switching. A negative bias of VTip = -5 V was written in a mirror opposite line and re-versed polarity (not shown). Overall, the PFM images show clear hysteresis behavior for all three films and the ability to switch polarization with a bi-ased tip.

Figure 4.7 Vertical PFM image for BTMNO-25 afterapplying a VTip +5 V over a 3 x 3 µm2 area.

59

5 Halide Perovskites

In this chapter, the results of the perovskite halide compounds are highlight-ed. The following chapter is based on results in Papers VI-X. The first por-tion of this chapter focuses on perovskite halides with Pb as the post-transition metal cation and their low-dimensional equivalents. The subse-quent section addresses the lead-alternative based halides and their electronic structure. And lastly, we report on nanocrystal halide perovskites used in a high-efficiency photovoltaic device.

5.1 Low-Dimensional Halides

Low-dimensional perovskite halides have emerged in the past few years largely catalyzed by the interest in stable halide solar cell materials. Their impressive efficiencies and remarkable stability relative to their 3D counter-parts (CH3NH3PbI3) has put this subclass of materials into broader focus and hold promise in energy-related applications123–125. The primary interest is on the x-ray spectroscopy, and how the reduced dimensionality effects the elec-tronic and band structures and it includes attempts to relate the spectroscopic findings to the material properties. In Paper VI, we report on the role of synthetic routes that form two com-pletely different structures. The one-dimensional (1D) structure, (BdA)Pb2I6 (BdA: butyl-1,4-diammonium), forms a one-dimensional inorganic chain with a disordered intercalation of the organic bications. The large cation causes the connectivity of the inorganic units to be lost in one or more di-mension in the perovskite network. The byproduct of a different synthetic route produce a two-dimensional structure in which (BdA)PbI4 (001) struc-tural planes to consist of octahedral corner sharing lead-iodide planes and organic bications separating the inorganic layers. The varying bond distances and angles are known to influence the orbitals that contribute to bonding of these materials126, and therefore the electronic structure can bear information to these structural changes.

60

Figure 5.1(a) shows the HAXPES spectra and x-ray absorption and emission spectroscopy on the two compounds deposited on a mesoporous TiO2 thin film substrate. The inset in Fig. 5.1(a) shows slight differences in the valence band edge position for the 1D and 2D material compared to the precursor PbI2, with a small shift for the 2D-(BdA)Pb2I6 towards smaller binding ener-gies. It was shown in Paper VI that the Pb 4f core-level shifts between the layered compounds are small (0.05 eV), however for I 3d, the core-level shift for (BdA)Pb2I6 and (BdA)PbI4 is about 0.12 eV. This data gives evi-dence that the positions of the valence band position shift largey due to io-dine p orbitals. In order to map the I p projected density of states, results from element-specific HERFD-XAS and RIXS at the I L2 edge is shown in Fig. 5.1(b,c). The absorption spectrum at the I L1 edge is a reflection of the unoccupied p density of states and coordination. The two data sets clearly show a similar spectral profile, but vary in intensity and energy position for the different structures, a change in the local structure. RIXS spectra in Fig 5.1(c) show a shoulder feature on the higher energy side of the spectrum that can be linked to changes in the hybridization with Pb 6s states. The Pb-I-Pb bond length for the 1D compound is 6.08 Å as compared to 6.16 Å for the 2D compound. These small changes to the local structure can affect the orbital overlaps of Pb and I states, and thus reflect in small changes to the electronic structure. In low-dimensional perovskites, the organic cation ionic radii and length is associated with varying structural properties that result in the tunability of material properties such as bandgap and excitons binding energies124,127–130. The next study is extended to tuning the inter-layer separation distance be-tween the single PbI6 octahedral layer (n = 1), as highlighted in Paper VII. The organic cation length is varied between the ~4 Å (BdA, butyl-1,4-diammonium), 6Å (HdA, hexyl-1,6-diammonium), and 8 Å (OdA, octyl-1,8-diammonium) and the changes to the electronic structure is surveyed through different x-ray spectroscopies. In contrast to the layered PbI2, the observed

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Figure 5.1 (a) Valence levels of the TiO2 film, PbI2 and (BdA)Pb2I6 and (BdA)PbI4measured with a photon energy of 4000 eV. (b) Plots of the unoccupied np states ofiodine in the conduction band and (c) RIXS spectra for (BdA)Pb2I6, (BdA)PbI4, and PbI2.

61

structural changes in the 2D hybrid materials results in a dimensional con-straint of the inorganic layer. The narrowing of VB states is demonstrated by a closer inspection of the valence band profile as shown in Fig. 5.2(a). This is an indication that the organic cation-linked 2D perovskites constitute more localized bands than PbI2 and MAPbI3. The valence bandwidth narrows as we go from the layer-separated PbI2 to the 3D MAPbI3, and lastly to the interlayer-separated compounds.

To evaluate iodine’s role in modifying the frontier valence states, we exploit the element-specific nature of HERFD-XANES and RIXS at the I L1 edge to probe occupied and unoccupied states as shown in Fig. 5.2. The structural changes are evident from the lineshape of the I L1 absorption spectra shown in Fig 5.2(b). An increase in intensity implies a larger number of unoccupied states for (BdA)PbI4 at regions marked I and II in the spectra. This increase may be the consequence of lower symmetry compared to (HdA)PbI4 and (OdA)PbI4, which is seen by minor changes to lattice positions as well as octahedral tilt. The change in interlayer distance also modifies the bandgap of the materials, with 2.37 eV for the BdA compound and 2.55 eV for the HdA perovskite. This effect can be seen in shifting of I 5p states toward lower energy loss for (BdA)PbI4 followed by (HdA)PbI4 and (OdA)PbI4. This trend follows the sequence of increasing band gap from (BdA)PbI4 (2.37 eV), (HdA)PbI4 (2.44 eV), and (OdA)PbI4 (2.55 eV) and is in agree-ment with the photoemission spectra reported earlier. In order to support the experimental evidence for band narrowing as result reduced dimensionality, DFT calculations based on Generalize Gradient Approximation (GGA) is presented in Fig. 5.3. The calculations highlights the contribution of Pb 6s and I 5p states in the VB and CB edges. As is clear-ly seen in Fig 5.3(a), the band narrows as the materials adopts a reduced dimensional structure as compared with PbI2 or MAPbI3. This is particularly

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Figure 5.2 (a) The VB photoemission spectra collected for all compounds at 4000eV. (b) HERFD-XAS at the I L1 edge that is projection of unoccupied I p states, and(c) RIXS spectrum plotted on an energy-loss scale.

62

evident for I states. The interactions between the interlayers show a direc-tional dependence, where I p states are mostly affected in the axial direction as shown in Fig 5.3(b). This demonstrates the orbital confinement effects in lower dimensional perovskite halides.

The band structures also show a directional dependence on the effect mass of charge carriers for all three layered compounds. A small carrier mass is ob-served for the in-plane direction within the inorganic planes and much larger perpendicular to the inorganic plane, shown in Fig. 5.4. This is in line with the expectation that the charge transport mainly occurs within the inorganic network131,132. In addition, a band dispersion for the direction perpendicular to the inorganic network is observed for the shortest alkyl chain (BdA)PbI4 perovskite (Figure 5.4(a)), showing an interaction and possible charge transport between the layers. This dispersion becomes flatter for (HdA)PbI4 (Fig. 5.4(b)) and disappears for the (OdA)PbI4 (Fig. 5.4(c)). This reveals that for distance longer than 6Å limit, there is no interaction between inorganic interlayers.

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(BdA)PbI4

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C H N

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MAPbI3 I Pb Total

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Figure 5.3 (a) Calculated total and partial density of states (DOS) of 2Dperovskites plotted alongside PbI2 and 3D MAPbI3. (b) Partial and totalDOS for iodine in (HdA)PbI4 showing the contribution to from both axialand equatorial iodine atoms. The effect of dimensionality clearly has aneffect on the orbital narrowing for iodine in the axial direction.

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5.2 Lead-Alternative Halide Perovskites

There is an ongoing effort to find new materials that can substitute lead with properties similar to that of lead-based compounds. The commonly used lead-based perovskite are associated with toxicity and concerns about ambi-ent stability15–17, and necessitating the pursuit of alternatives beyond the lead-based structures. The choice of a divalent cation not only determines the crystal unit structure, but also tunes the optical properties of the resultant material. Several studies suggest the Pb2+ cation plays a definitive role in determining the electronic structure of Pb-based perovskites, particularly at the valance and conduction band edges23,133. Metal substitutions by Sn2+, Sb3+, and Bi3+ are most suitable due their similar ionic radii and valence elec-tron configuration, ns2. The incorporation of these materials in solar cell devices is still at its nascent stage, however have demonstrated good results that merit further investigation134. The electronic structure of the Pb-based perovskites maintains a stable Pb2+ oxidation state that results in a mostly 3D crystal structure. The different valency and ionic radii imposed by a substituted cation imposes steric hin-drance that can trigger the formation of a lower symmetry structure or lay-ered systems. Regardless, it is likely that the principal building block in the crystal structure is a MX6 octahedra that can be joined in corner-sharing, edge-sharing, or vertex-sharing and repeated along a crystallographic direc-tion. It is well established that the A-site cation determines the structure, stability, and bandgap of the material. However, it is preferentially ionic and has shown to have minor influence on the band edge electronic structure.

Figure 5.4 Band structure plots for (a) (BdA)PbI4, (b) (HdA)PbI4, and (c) (OdA)PbI4together with K-space paths with the first Brillouin zone and their estimated bandgap values calculated with DFT.

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Therefore, the electronic structure information obtained through various experimental and theoretical methodologies can serve as an appropriate comparison against the best performing HP material. In Paper VIII, x-ray spectroscopy results on bismuth-based halide materials (Cs3Bi2I9 and BiI3) are discussed. Bismuth-based hybrid perovskites also bear similarities to the Pb-based analogues that also adopt a stable d10s2p0 electronic configuration with covalent Bi 6p/6s and I 5p at the top of the valence band and mostly Bi 6p at the bottom of the conduction band135. It was suggested that this specific orbital occupancy leads to anti-bonding states that form defect states that largely lie within the valence band, a de-sired attribute to reduce carrier recombination53,135,136. The experimental x-ray results are compared against DFT calculations. We first probe the bulk valence band structure for the two bismuth com-pounds and the related precursor CsI using HAXPES as shown in Fig. 5.5. The overall lineshape of Cs3Bi2I9 has similarities to the precursor, with mod-ification in the valence band edge position in between the low bandgap BiI3 and high bandgap CsI (BiI3: 1.67 eV, Cs3Bi2I9: 2.2 eV, CsI: 6.3 eV). The doublet peak near ~ 12 eV corresponds to Cs 5p orbitals, states that do not participate in bonding as expected. The calculated partial and total density of states (DOS) for both Cs3Bi2I9 and BiI3 alongside the HAXPES data is pre-sented in Fig. 5.5 (b). The DFT results are in close agreement to the experi-mental data. The I p states have substantial weight at lower binding energies of the valence band, and hybridize with Bi 6s lone pair states towards the edge of the valence band.

Figure 5.5 (a) Valence band spectra for the three samples at 4000 eV (b) calculatedand experimental VB for Cs3Bi2I9 and BiI3 along with experimental spectra.

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HERFD-XANES and RIXS taken at the I L1 edge are shown in Fig 5.6 along with the projected DOS using the Local Density Approximation (LDA) with Spin-Orbit Coupling (SOC). All three samples show distinct spectral charac-teristics due to the different coordination and displacement of iodine atoms in the unit cell. The absorption edge, defined at the maximum of the second derivative of the absorption spectrum, shifts to higher energy for Cs3Bi2I9 and CsI and correlates with the concomitant increase in bandgap (see Fig. 5.6(a)). The RIXS presented in Fig. 5.6(b) on an energy loss scale is charac-teristic of occupied I 5p states. The energy difference between the elastic peak and the onset of the inelastic structures on the energy loss scale corre-lates with the size of the band gap in these compounds. Therefore, the shift of such an onset to a lower energy loss by 0.4 eV in BiI3 as compared with Cs3Bi2I9 indicates a small band gap in the former. This information can pro-vide insight into the changes of I p states, and how they ultimately affect the band gap.

We now extend the study of bismuth compounds with the commonly used cation methylammonium (MA) at the A-site of the A3M2X9 structure as out-

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Figure 5.6 (a) HERFD-XAS taken at the I L1 edge for all three samples.(b) RIXS spectrum plotted on an energy loss scale. The projected unoc-cupied density of states for (c) Cs3Bi2I9 and (d) BiI3.

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lined in Paper IX. The large cation and the mismatch in ionic radii force the structure to adopt a HP-like structure that results in a (CH3NH3)Bi2I9 com-pound, in a iododi-bismuthate crystal structure (A3Bi2I9). This material has exhibited promising optoelectronic properties, with a photovoltaic efficiency of nearly 4% based on a optimized thin-film deposition technique137. The vapor assisted solution process (VASP) at different reaction times yielded tunable optoelectronic parameters, a highly uniform and compact film, im-proved stoichiometry, reduced metallic content, and desirable absorbing properties138. HAXPES was used to investigate the bulk electronic levels of MA3Bi2I9 (MA: CH3NH3) synthesized at various reaction times in the VASP process. The two precursors, MAI and BiI3 as subjected to varying annealing times after spin coating. The transition from the two precursors to the desired ma-terial can be monitored using the core-level photoelectron signals. An ap-proximate estimation of the material stoichiometries can be made using the relative intensities of the experimental core levels and taken differential pho-toionization cross-sections into account139. The stoichiometry estimation for I/Bi is listed in Table 1, with values reasonably close to the expected value of 4.5. This stoichiometry points to marginal iodine deficiencies within the crystal, a common bulk defect found in many organometallic halides140.

Table 5.1 HAXPES stoichiometry data for the measured samples.

The experimental Bi 4f and valence band signals are presented in Fig. 5.7(a). The main contribution of the Bi 4f single is around 159 eV, indicating a Bi3+ oxidation state141. The shoulder peaks at lower binding energy (~157 eV) is metallic Bi0, which is likely due to the decomposition organic metal halides and BiI3.142 As the reaction time is increased, the detrimental metallic Bi0 content is reduced, concurrent with photovoltaic efficiencies and improved film quality, demonstrating the effectiveness of longer vapor exposure times. Although the MA cation is much larger than the Cs atom studied previously, the optical transition is expected to be the same143, as shown in the valence band profile in Fig. 5.7(b). The dashed red line (see inset Fig. 5.7(b)) indi-cates where the linearly extrapolated experimental spectra intersect with the baseline, an estimated measure of the valence band edge position. The VB offset are approximately the same for the entire set, except for reaction time of 5 min, which has a substantially higher metallic bismuth concentration. The energy position of the valence band maximum lie mostly in line with

Sample BiI3 BiI3 + MAI 5 min

BiI3 + MAI 15 min

BiI3 + MAI 25 min

BiI3 + MAI 45 min

I 4d/Bi5d 3.03 4.31 4.30 4.50 4.28 I 3d/N1s - 2.94 2.95 2.82 2.46

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BiI3, indicating the corner-sharing counter cation CH3NH3+ seems not rele-

vant to the density of states in the valence band edge. The features between 3-5 eV is composed of I 5p and Bi 6p and lone pair 6s2 anti-bonding states, similar to that of the Cs3Bi3I9 compound discussed in the previous section.

In the previous sections, lead-alternative halides with Bi at the post-transition metal cations compounds were investigated in terms of their elec-tronic structure. There are several other compounds that are related to this topic, however the results are not compiled into a manuscript. Regardless, it is worth highlighting some of these preliminary results. A recent HAXPES study was conducted on three post-transition metal HPs all with a lone pair valence configuration (Sb3+: 5s2, Pb2+: 6s2, and Bi3+: 6s2). The cation at the A-site was varied between the smaller inorganic Cs and the larger and organ-ic methylammonium (MA). The valence band photoemission spectrums are shown in Fig. 5.8 and are normalized to peak maximum near 3.5 eV for all samples. The larger cation MA+ pushes the valence band edge closer to the Fermi level for all three post-transition metal compounds, suggesting the materials may become more n-type. The reported bandgap values from pre-vious studies have shown to be smaller for MA to that of Cs-based hal-ides141. This suggests that the cation size plays a definite role in alternating the bandgap as well the position of the valence band maximum (VBM). The Sb3+ HPs show a pronounced structure near 1.3 eV, which has been showed to belong to lone-pair 5s2 states based on band structure calculations144,145. The peculiarities of the electronic structure and the projected density of states near the band edges need further investigation through complimentary spectroscopy techniques, the intentions of future work.

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Bi 4f

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Figure 5.7 (a) Core-level peak for Bi 4f showing a decrease in metallic content withdeposition times. (b) Valence band spectrum for all samples and the annealed BiI3 film and BiI3 + MAI(v) for different reaction time.

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5.4 Nanocrystal Halides

Along with the rapid development of bulk perovskite deposited with solution process techniques, the research into colloidal perovskite nanocrystals (CPNs) has also emerged very recently with favorable results146. The poten-tial for CPNs are enormous, as colloidal quantum dots (CQDs) have already been building block components to light emitting devices (LEDs), photode-tectors, and solar cells56,147. We recently reported on a one of the best-performing CPNs-based solar cell using CsPbI3 and PbS as functional units in the light-harvesting layer of the photovoltaic devices55. The material is based on CsPbI3 CQDs grown directly on top of an optically active PbS CQD – an architecture that is coined ‘core-shell quantum dot structure’ that results in an observable quantum-confinement effect148. This work is report-ed in Paper X and only the relevant results will be discussed here. Figure 5.9 displays the core electronic levels and the valence band for three systems: oleic-acid capped PbS (PbS-OA), PbI2 shell and PbS core CQD (PbS-PbI2), and the HP CsPbI3 grown on a core PbS (PbS-CsPbI3). The core-levels (Pb 4f, Cs 4d, I4d) are normalized to the S 2p extending from the PbS underlayer. Based on the S 2p ratio with respect to Pb 4f core-level signals

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Figure 5.8 HAXPES VB spectra for HPs with three different B-site metal: Sb3+,Bi3+, and Pb2+. The VBM shifts to lower B.E. as a result of the larger cation.

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for the PbS-OA CQD sample, an approximation of the shell materials stoi-chiometry can be made. Comparing the additional Pb 4f signal in PbS-CsPbI3-P to the intensity of the I 4d and using theoretical photoionization cross-sections149, we obtain a stoichiometric ratio of I/Pb = 3.1 and I/Cs = 2.7, agreeing with the expected stoichiometry. The relative intensities of the Pb 4f signal, when compared against the Pb 4f intensity from PbS-OA sample, points to a very thin surface layer for PbI2 and much stronger for CsPbI3. Only minor core-level shifts in the S 2p and Pb 4f signal of PbS-CsPbI3 (~ 0.15 and 0.06 eV respectively) are observed, suggesting that no large dipole layer exist at the interface of PbS/CsPbI3 CQDs. In addition, the relative position of the valence band maximum (indicated by dashed lines) is shifted to lower binding energy for PbS-CsPbI3 by ≈ 0.4 eV versus the Fermi level in comparison to the other two samples. Since the capping CsPbI3 layer contributes significantly to the overall measured signal, the edge shift is therefore likely from the capping CsPbI3 layer. This suggest that the VB edge of CsPbI3 is 0.4 eV further away from the Fermi level that that of PbS-OA, which together with the absorption properties and work function meas-urements (see Paper X) indicate electronic energy levels that favor charges carrier confinement.

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Pb 4f Cs 4d I 4d VB

Figure 5.9 Core level x-ray photoelectron spectra normalized to S 2p signal of Pb4f, Cs 4d, I 4d for the PbS-PbI2 and PbS-CsPbI3-P based CQD solids, respective-ly, and valence band spectra (normalized to maximum intensity).

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6 Outlook

This thesis presented the versatility of two classes of compounds through x-ray spectroscopic and thin film characterization. A number of complimen-tary techniques along with different levels of theory were applied in order to describe the electronic structure and relate the spectroscopic information to their macroscopic properties. In order to develop a more detailed picture of LaFe1-xMoxO3 unusual ground state antiferromagnetism with metallicity could be through Kβ x-ray emission spectroscopy and Fe K edge RIXS. A more complex understanding on the charge disproportionation could be understood by describing the nominal spin state for the two separate Fe species in the metallic compounds. This could further validate the disproportionation observed at the Fe site. Addi-tionally, fabrication of thin films would be the obvious next step to definitely validate metallicity at low temperatures without the complications from grain boundary effects that smear the inherent bulk properties of the materi-als. The most promising outlook in the entire series of compounds studied in this thesis is perhaps in the development of BaTi1-x(Mn1/2Nb1/2)xO3 ferroelectrics. Thin films have already been fabricated and characterized with promising results. The next step would be to incorporate the films into photovoltaic devices to determine if the inherent polarization and the reduced bandgap can result in efficient devices. For a more fundamental understanding, the role of d0 Nb5+ is still unclear other than its charge-compensating role in the unit cell. EXAFS studies at the Nb K edge could determine if this dopant is also off-center at the B site, and whether other d0 compounds could perhaps even increase the remnant polarization as a function of doping. These results could offer new strategies in doping ferroelectrics to achieve new functional-ities in electronic devices and to energy-related applications.

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7 Acknowledgements

I would like to start first by thanking my three advisors Prof. Dr. Olof Karis, Prof. Dr. Håkan Rensmo, and Dr. Sergei Butorin. Thank you for your guid-ance, patience, and mostly your support in giving me the freedom to do many different projects. I am grateful for your guidance in x-ray spectrosco-py and material physics. I of course must extend my sincere thanks and gratitude to Dr. Somnath Jana and Dr. Bertrand Philippe as my ‘day-to-day’ advisors, whom were instru-mental in guiding me in the principles of halide perovskite materials, oxide materials, and x-ray spectroscopy. Your support is truly appreciated.

I also extend my thanks to colleagues Shyam and Soham from the Indian Institute of Science whom provided me with very interesting samples and valuable discussions. I hope that we can continue to develop the BTMNO project further and to finally realize our goals of incorporating the material into functioning devices. Tremendous thanks to Dr. Soham Mukherjee for helping and guiding me through the grind and grit in the last year.

The thin film BTO project would not have been possible without the support of Henri Frederiksen and Dr. Alexei Kalaboukhov at Chalmers University. Thanks for all your incredible assistance in realizing this project. I am in-debted to you for introducing me to Pulsed Laser Deposition and thin film characterization. Much of the perovskite halide studies would not be possible without support from colleagues at KTH. Thank you Majid and Dr. Gardner for involving me in on your projects early on in the PhD. I must also thank Sagar Jain, Malin B. Johansson, and Xiaoliang Zhang who have given me many unique sam-ples to study and wonderful discussions, and there is still more work to be done. It was a great pleasure working with Dr. Tomas Edvinsson and Dr. Meysam Pazoki. I think we make a good team, and hope we can find the time and resources to do all the projects we’ve discussed over the years into the fu-ture. Thanks for all your support.

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I must also thank Dr. Ute Cappel for her wonderful support in photoelectron spectroscopy. I am always learning more effective ways to be a thoughtful and efficient experimentalist at beamtimes with you, and hope to be on more with you in the future. Dr. Venkata has also give me a wonderful opportunity to work with fabrication and thin film devices, and I hope to realize our dream projects soon. Thanks very much for the involving me in your team and on some very interesting and difficult projects.

My colleagues at the Division, thank you for sharing in wonderful discus-sions, laughs, and your own stories over these years. I have enjoyed working with you and hope to keep the good vibes going.

To my dear family, who have always encouraged and believed in my whirl-wind journey in the last few years. You are the core unit in my life. And lastly, thank you Camilla for sharing this journey with me.

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8 Populärvetenskaplig Sammanfattning

Behovet av alternativa energikällor blir för varje år som går en allt mer akut fråga. Trots att kostnaderna för kiselbaserade solceller för närvarande är rekordlåga finns det betydande hinder i form av tillverkningskostnader för kisel och andra tunnfilmstekniker som kan göra det allt svårare att tillräckligt minska kostnaderna för att ge en stark ekonomisk drivkraft mot förnyelsebar energi. Liknande problem finns för den mikroelektronik som är nödvändig för att möjliggöra en fortsatt utveckling av informationsteknologin, som närmar sig skalbarhetsgränsen för den nuvarande CMOS-tekniken, trots att en stor energianvändning fortfarande krävs. Vårt beroende av dessa två tek-niker kräver en inriktning mot nya material som kan möta de behov vi ser i framtiden. Det finns nya teknologier som väcker stort intresse hos forskare inom sol-cellsfältet där nya material används som är effektiva på att uppta solenergi och dessutom till en bråkdel av kostnaden för dagen teknologi. Detta nya fält använder halidperovskiter som är kända för sin höga ljusabsorption och goda opto-elektroniska egenskaper som kan skräddarsys för specifika ändamål. Den mest framträdande egenskapen är tillverkningsmetoden som sker i lös-ningar som varken kräver höga temperaturer, höga tryck eller dyrbar utrust-ning för tunnfilmsproduktion, vilket potentiellt kan göra produkter som base-ras på dessa material avsevärt billigare. Övergångsmetallernas oxider, i synnerhet oxider med så kallad perovskit-struktur ABO3, är en lovande materialklass som uppvisar en nästan obegrip-lig rikedom av egenskaper som kan användas för att skapa funktionalitet på komponentnivå. Det finns en stor tilltro till att dessa typer av material har potential att möjliggöra ytterligare ökning i beräkningstätheten och energief-fektiviteten hos kretsar. Utvecklingen inom beläggning av mycket tunna oxidfilmer, tillverkning av heterostrukturer och karaktäriseringsmetoder har fört oss in i en ny era där komplexa oxider kan användas för multifunktion-ella komponenter. Denna avhandling har studerat en speciell form av dopade ferroelektriska filmer som växts genom att deponera material under kontrol-lerad miljö med hjälp av intensiva pulsade lasrar. Detta system är särskilt intressant då det visar att det är möjligt att växa modifierade ferroelektriska material i tunnfilmsform med i stort bibehållna egenskaper med hjälp av denna metod. Att kunna åstadkomma ett ferroelektriskt material med ett

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mindre så kallat gap, vilket möjliggör att en större del av del tillgängliga energin i solljuset skulle kunna omvandlas till elektrisk energi, i form av en tunnfilm är mycket relevant för att kunna skapa nya solceller där ett av de grundläggande problemen i cellen, laddningsseparation, sköts av den inne-boende polarisationen i dessa material. För att förstå vad som ger upphov till de så eftertraktade egenskaperna hos dessa material som gör dem relevanta för användning inom solceller och komponenter av olika slag är det mycket värdefullt att kunna studera elektronstrukturen hos materialen i detalj. Denna avhandling försöker belysa hur de unika egenskaper som gör materialen användbara i en rad tillämp-ningar uppkommer, genom att beskriva materialens elektronstruktur så detal-jerat som möjligt. Den systematiska undersökningen av elektroniska egen-skaper hos övergångsmetalloxider and metallhalider är en central del för att förstå deras egenskaper. Den grundämnes- och platsspecifika naturen hos

Figure 8.1 (a) De många olika former som metallhalider med perovskitstruktur kan anta. (b) I oxidmaterial har det visat sig vara möjligt att koppla samman flera olika frihetsgrader vilket kan leda till att en egenskap kan kontrolleras genom att manipulera en annan. Ett eftersökt exempel är att styra det magnetiska tillståndet med en elektrisk spänning. Sådan funktionalitet gör dessa material intressanta för olika tillämpningar. (c) I ett ferroelektriskt material kan polarisat-ionen (fördelningen av positivt och negativt laddade joner) ändras på ett bestå-ende sätt med en pålagd spänning.

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flera röntgenbaserade spektroskopiska tekniker ges möjligheter för att få en komplett bild av de elektroniska egenskaperna hos materialen i fråga. Hos metallhaliderna undersöks valens- och ledningsbanden med röntgen-absorption, röntgenemission och fotoelektronspektroskopi. De spektrosko-piska resultaten identifierar vilka tillstånd som formar valensbandet och led-ningsband, vilka är kritiska parametrar att känna till för optoelektroniska komponenter. Därutöver har röntgenbaserad spektroskopi använts för att demonsterara den stereokemiska aktiviteten för vad som benämns lone-pair tillstånd (5s2 och 6s2) för flera olika halidföreningar och deras påverkan på de kemiska, strukturella och elektroniska egenskaperna hos materialet. Dess-sutom har positionerna för jodtillstånd och valensbandstillstånd i lågdimens-ionella bly-baserade komponenter undersökts i förhållande till sina tredi-mensionella motsvarigheter. Denna information belyser de intressanta materialegenskaperna och deras användning i vad som ofta benämns tredje generationens solcellsforskning. Detta avhandlingsarbete har också behandlat flera system av funktionella oxider. Speciellt har en ny grupp av metalloxider baserade på dopning av antiferromagnetiskt LaFeO3 , syntetiserats och studerats med hjälp av rönt-genabsorption, röntgenemission och fotoeletronspektroskopi. Elektronstruk-turen är komplex, och uppvisar både delokaliserade, eller så kallat itineranta tillstånd, såväl som lokaliserade delar som ger upphov till ett unikt tillstånd för ett material där antiferromagnetism, metallicitet och laddnings-disproportionering samexisterar. Med resultaten från resonanta fotoemission kan vi bestämma att Fe-tillstånden i de båda magnetiskt ordnade oxiderna inte bidrar till ledningsförmågan, medans molybdentillstånden uppvisar en intinerat karaktär vid Ferminivån. En undersökning utförd med resonant inelastisk röntgenspridning på LaFe1-xMoxO3 och dubbelperovskiten Sr2FeMoO6 uppvisar strukturer som vi har kunnat tolka som excitationer inom tillstånd av d-karaktär, s.k. d-d-excitationer, vilket visar på den starkt korrelerade naturen av dessa övergångsmetallmaterial.

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